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Processing and thermal properties of an Mo5Si3C–SiC ceramic
Intermetallics 16 (2008) 854–859
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Intermetallics journal homepage: www.elsevier.com/locate/intermet
Processing and thermal properties of an Mo5Si3C–SiC ceramic Andrew A. Buchheit a, Greg E. Hilmas b, *, William G. Fahrenholtz b, Douglas M. Deason c, Hsin Wang d a
Coorstek, Inc., Golden, CO 80401, USA Materials Science and Engineering Department, University of Missouri-Rolla, Rolla, MO 65409, USA c U.S. Army Space & Missile Defense Command, Building 5220, Redstone Arsenal, AL 35898, USA d Oak Ridge National Laboratory, High Temperature Materials Laboratory, Oak Ridge, TN 37831, USA b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 10 July 2007 Received in revised form 25 February 2008 Accepted 8 March 2008 Available online 3 June 2008
Composite ceramics containing w25 vol% of the Mo–Si–C ternary compound and SiC (referred to as MS) were reaction hot pressed up to 96% relative density. The electrical resistivity of the composite processed at 2150 C decreased from w4.60 U cm at room temperature to 4.15 U cm at 700 C. Thermal conductivity of the MS ceramics processed at 2150 C was above 100 W/m K at room temperature, decreasing to between 62 and 68 W/m K at 500 C. Using the Eucken model for thermal conductivity, the interconnected SiC phase in the MS materials was calculated to have a room temperature thermal conductivity between 160 and 170 W/m K. As compared to a baseline SiC composition, the continuous SiC in the MS materials had an average increase in thermal conductivity of w50% over the temperature range of room temperature to 500 C. This increase in thermal conductivity was attributed to the accommodation of impurities that would typically be present in SiC grains and grain boundaries (e.g., N and O) into the ternary phase. Ó 2008 Elsevier Ltd. All rights reserved.
Keywords: A. Molybdenum silicides B. Thermal properties C. Reaction synthesis F. Diffraction F. Electron microscopy, scanning
1. Introduction
is written as Mo5Si3C and crystallizes in the D88 structure (P63/ Stoichiometry of the ternary compound ranges from 50 to 60.7% molybdenum, 27.1 to 36.8% silicon and 8.3 to 15.6% carbon based on the Nowotny phase diagram and the eutectic temperature in the SiC–C-ternary phase field was reported to be w1900 C [8]. The composition Mo4.89Si3C0.686, which is within the ternary compound stoichiometry, has a coefficient of thermal expansion (a) anisotropy (ac-axis/aa-axis) of w2 with ac-axis having a room temperature value of 9.2 ppm/ C. The single crystal thermal conductivity of the ternary compound with the same stoichiometry was reported to be 8.5 W/m K in the ½1120 direction. A room temperature conductivity of w10 W/m K was measured for polycrystalline Mo5Si2.22C1.16, which is the composition stable in contact with C and MoC [9]. In contrast to the ternary compound, silicon has a thermal conductivity of w102 W/m K and an isotropic CTE of w4.4 ppm/ C from room temperature to 900 C [10,11]. As a result, Si melt infiltrated SiC has a thermal conductivity between 100 and 175 W/m K at room temperature [10]. With these differences in property values between the ternary Mo–Si–C phase and Si, the thermal conductivity of the SiC–Mo5Si3C ceramics is expected to be lower than that of Si-infiltrated SiC ceramics. The objective of this paper is to characterize the thermal properties of SiC–Mo5Si3C ceramics in addition to the mechanism of ternary phase formation. The in situ reaction and densification process, resulting in an overall composition in the range of those studied by Zhu et al., were explored as an alternative to melt infiltration. mcm).
The properties of silicon carbide ceramics depend on the processing method. Melt infiltration is one of the methods employed to manufacture SiC ceramics. Generally, the infiltrant is silicon, which limits the use of melt infiltrated materials to temperatures below the melting temperature of silicon (w1400 C). Average strengths of Si-infiltrated SiC are reported to be w350 MPa at room temperature and 200 MPa at 1400 C [1]. Within the past decade, Zhu et al. studied the ternary Mo–Si–C compound as an alternative infiltrant [2] due to its high melting temperature (w2100 C) and its ductile to brittle transition around 1300 C [3–5]. To date, mechanical properties have been measured for compositions containing 70 vol% SiC with 25–30 vol% of the ternary compound [5,6]. From the latter studies, approximately 90% of the room temperature strength (290 MPa for 92% dense materials) was retained up to 1700 C [7] and fracture toughness increased up to w8 MPa m1/2 above the ductile to brittle transition temperature of the ternary compound [6]. The properties of the ternary Mo–Si–C compound are significantly different than those of silicon. The stoichiometric compound
* Corresponding author. Materials Science and Engineering Department, University of Missouri-Rolla, 225 McNutt Hall, Rolla, MO 65401, USA. Tel.: þ1 573 341 7205. E-mail addresses: [email protected] (A.A. Buchheit), [email protected] (G.E. Hilmas). 0966-9795/$ – see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2008.03.010
A.A. Buchheit et al. / Intermetallics 16 (2008) 854–859
2. Experimental procedure 2.1. Powder processing and specimen fabrication Molybdenum powder was 99.95% pure and had an average particle size in the range of 2–4 mm (M-1109, Cerac, Inc., Milwaukee, WI). Silicon carbide powder was >95% b-SiC (BF-12, H.C. Starck, Goslar, Germany) with a reported average particle size of 0.6 mm and a surface area of 11.0–13.0 m2/g. The reactant mixture contained 15 vol% molybdenum with the remainder being SiC. Powders were attrition milled (Model 01-HD, Union Process, Akron, OH) for 1 h in hexane at 600 rpm using a fluoro-polymer coated spindle, 750 cc bucket and 3/16’’ (4.76 mm) diameter SiC media. Powders ¨ CHI Laborwere dried by rotary evaporation (Rotavapor R-124, BU technik AG, Flawil, Switzerland) at 70 C, 75 rpm and 40 Pa vacuum. Specimens will be referred to as MS with a numerical prefix designating the processing pressure (in MPa) and a numerical suffix designating the processing temperature (in C), as shown in Table 1. The batch has an overall composition that falls in a three phase field containing SiC, C, and the ternary compound. Several assumptions were made when predicting the stoichiometry of the ternary compound from the atomic percentages of Mo, Si and C in the batch. The first assumption was that molybdenum and silicon in the ternary compound only resided on Mo and Si lattice sites, and the maximum number of sites in the unit cell was 5 and 3, respectively. Carbon was assumed to occupy all of the C sites, but also substitute onto Si sites due to its similar bonding characteristics and smaller atomic size compared to Si. This gives the constraint that C þ Si must be 4. Following these rules and using the Nowotny Mo–Si–C phase diagram [8] the stoichiometry of the ternary compound in equilibrium with SiC and C was predicted to be 53.81% Mo, 30.60% Si and 15.59%, which corresponds to Mo4.66Si2.65C1.35. A graphite die was lined with graphite foil (UCAR Carbon Company, Inc., Columbia, TN). Powder was added to the die and it was pressed uniaxially at 15 MPa (Model C, Carver, Inc., Wabash, IN). Specimens were processed in a graphite element hot press (HP20-3060, Thermal Technology, Inc., Santa Rosa, CA). During the hot pressing cycle, vacuum was maintained at <16 Pa up to 1950 C and then the furnace was backfilled with helium. Above 1950 C, a flowing He atmosphere was maintained at a slightly positive pressure (w105 Pa). At 1100 C, a uniaxial pressure of 25 MPa was applied and the pressure was increased to either 40 or 50 MPa at 1900 C, where temperature was held for 30 min. The final hold was 1 h at either 2000 or 2150 C. Cooling occurred at w30 /min until 1500 C. The pressure was released when the furnace temperature reached 1750 C. Below 1500 C, the furnace was allowed to cool uncontrolled. Three different combinations of processing conditions were used to fabricate specimens (Table 1). Bulk densities were determined geometrically using disks w12.5 mm in diameter and w3 mm thick. For comparison, a baseline alpha SiC (UF-10, H.C. Starck, Goslar, Germany) composition containing 1 wt% Al and 3 wt% C was hot pressed at 2000 C and 32 MPa for 1 h. 2.2. XRD and SEM analyses Phase evolution was investigated by X-ray diffraction analysis (XRD; Scintag XDS 2000, Thermo Optek Corporation, Franklin,
Table 1 Specimen designations and the corresponding processing conditions Designation
Pressure (MPa)
Temperature ( C)
50MS2000 40MS2150 50MS2150 SiC (1 wt% Al)
50 40 50 32
2000 2150 2150 2000
855
MA) on powder samples that were heated in a graphite furnace (Model 1000-3060-FP20, Thermal Technology, Inc., Santa Rosa, CA) with the same temperature and atmosphere profile as those hot pressed, except that temperature did not exceed 1900 C. The final composition of the hot pressed (HP) specimens was determined using XRD analysis of powders produced by crushing and grinding hot pressed materials. Powders were scanned from 20 to 80 2q using Cu Ka radiation. Patterns were indexed to powder diffraction file cards for 3C SiC (29-1129), 6H SiC (29-1131), Mo2C (35-0787), Mo5Si3 (34-0371) and Mo4.8Si3C0.6 (43-1199). Specimens were sectioned perpendicular to the hot pressing direction and polished to 3 mm for examination by scanning electron microscopy (SEM; S-570 SEM, Hitachi, Tokyo, Japan). Multiple images of polished surfaces were examined using image analysis software (ImageJ, National Institute of Health, Bethesda, MD) to estimate volume percentages of phases present in the hot pressed specimens.
2.3. Thermal testing Electrical resistance was measured using a four point probe potientiodynamic setup (Solartron 1470, Solartron Analytical, Farnborough, Hampshire, England). Bars w20 mm 1.5 mm 2 mm were notched in four places around their perimeters to allow attachment of platinum electrode wires. The wires were wrapped around the specimens and silver paste was applied to improve contact. Specimens were heated to 700 C under flowing argon to sinter the silver paste. During cooling, resistance measurements were recorded every 50 C between 700 and 400 C and then every 100 C to room temperature. Electrical resistivity was calculated using the following equation:
r ¼
RA L
(1)
where r is resistivity, R is the measured resistance, A is the crosssectional area of the specimen and L is the gauge length between the two voltage (inner) leads. Heat capacity was measured every 50 C between room temperature and 500 C in flowing high-purity argon with a ramp rate of 20 C/min (STA 409C, Netzsch, Selb/Bavaria Germany). A sapphire disk was used as a standard and all measurements were taken during the heating portion of the cycle. Data reported are from specimens w5 mm in diameter and <1 mm thickness. The thermal diffusivities of 50MS2000, 40MS2150 and the baseline SiC were measured at Oak Ridge National Laboratory (ORNL) in the High Temperature Materials Laboratory (HTML) user facility using a commercial laser flash system (Flashline 5000, Anter Corp., Pittsburgh, PA) on w12.5 mm diameter disks of w3 mm thickness. The reported thermal diffusivities were the average of three measurements made every w100 C and based on the Clark and Taylor data (ASTM E1461-01). Two separate furnaces with temperature ranges of RT to 500 C and 500 to >1900 C were used to determine the diffusivity to 1000 C. A silicon-based infrared pyrometer was used for temperature measurements. The thermal diffusivities of specimens of compositions 40MS2150 and 50MS2150, the same size as those tested at ORNL, were also measured every 100 C from room temperature to 900 C (LFA 457, Netzsch, Selb/Bavaria Germany) using a system equipped with an InSb detector. Thermal conductivities were then calculated using Eq. (2) where d is the thermal diffusivity, r is the bulk density and CP is the heat capacity. k ¼ drCP
(2)
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3. Results and discussion 3.1. Phase evolution The phase evolution study revealed molybdenum and SiC did not react below 700 C (Fig. 1). At 900 and 1100 C, Mo5Si3, Mo2C and a trace amount of Mo5Si3C1 (at 1100 C only) were detected along with SiC. A study by van Loo et al. [12] on the Mo–Si–C system reported the formation of Mo5Si3, Mo2C and Mo5Si3C at 1200 C, which agrees well with the analysis at 1100 C in the present study. At 1300 C, both Mo5Si3C1 and Mo2C were detected. Above 1750 C, only SiC and Mo5Si3C1 were present. Literature regarding SiC–Mo diffusion couples found the formation of both Mo2C and Mo5Si3 by 1200 C and that Mo5Si3 converted to Mo5Si3C1 between 1400 and 1500 C [13]. A previous study of SiC–Mo diffusion couples reported that only Mo2C and Mo3Si2 (Mo5Si3) formed when the couple was held for 9 h at 1700 C [14]. Based on the electron microprobe analysis performed by Morozumi et al. and the Nowotny Mo–Si–C phase diagram [8], that specimen had a layered structure with, most likely, a layer of Mo5Si3C1 in contact with SiC, a layer containing a mixture of Mo5Si3 and Mo2C, and then a layer of Mo2C or MoC1x in contact with Mo. The phase evolution study revealed that between the temperature ranges of the van Loo and Nowotny phase diagrams (1200–1600 C) Mo5Si3 and SiC were no longer in equilibrium. This was shown to occur by 1300 C with only a 30-min hold in the present study as no Mo5Si3 was detected by XRD. Also, above 1750 C, no detectable amount of Mo2C was found. Rietveld analysis of the XRD patterns, collected as part of the present study for powders heated to temperatures from 1300 to 1900 C (Fig. 1), was used to determine the lattice parameters of the ternary compound with a stoichiometry of Mo4.66Si2.65C1.35 (Table 2). Different stoichiometries of the ternary phase have been examined and lattice parameters have been reported in the literature [15]. Reported lattice parameters for Mo4.7Si3C0.6 hot pressed at 1500 C are 7.290 Å for the a and b axes and 5.043 Å for the c-axis [15]. The lattice parameters have also been reported for the ternary compound without specific stoichiometries, although the ternary compound was in equilibrium with either Mo2C (ranges from Mo5Si2.19C1.05 to Mo5Si2.66C0.69) or MoSi2 (ranges from Mo3.93Si2.96C1.04 to Mo4.31Si3C0.69), which will be referred to as Mo-rich and Mo-poor regions, respectively. The a-axis lattice parameter for the Mo-rich ternary compound was 7.2878 Å and the c-axis parameter was 5.5069 Å. For the Mo-poor compound, the a-axis parameter was 7.2884 Å and the c-axis parameter was 5.0164 Å [15]. The c-axis parameter calculated in the present study falls within previous ranges reported in the literature. Notice the
Fig. 1. Phase evolution in Mo–SiC mixtures between 700 and 1900 C with selected phase-specific peaks labeled.
Table 2 Lattice parameters and crystal volume of Mo4.66Si2.65C1.35 from the phase evolution study at their corresponding temperatures Temperature
a/b axis (Å)
c-axis (Å)
c/a
Volume (Å3)
1300 1500 1750 1900
7.26(4) 7.26(8) 7.26(7) 7.26(8)
5.04(7) 5.04(4) 5.03(8) 5.03(9)
0.695 0.694 0.693 0.693
230.67 230.73 230.43 230.60
Mo stoichiometry is nearly the same as that reported by Suzuki et al. (Mo4.66 vs. Mo4.7) and the average measured c-axis was 5.042 0.004 Å, suggesting the cell height was controlled by the Mo–Si bonds. However, the lower a-axis parameter was attributed to the substitution of C onto Si sites in combination with fully occupied C sites, which should result in a decrease in the size of the Mo–Si octahedra. The lattice parameter analysis was performed to allow the theoretical density to be calculated. Using the lattice parameters, the crystal volume of Mo4.66Si2.65C1.35 was calculated to be 230.61 0.13 Å3 with a molar weight of 529.58 g/ mol. These values result in a theoretical density of 7.624 g/cm3 for Mo4.66Si2.65C1.35 as compared to 7.89 g/cm3 as reported in PDF card 43-1199 for Mo4.8Si3C0.6. The decrease in Mo content and the increase in C content of the material examined in this study led to a slightly lower theoretical density despite the decrease in unit cell volume. Several chemical expressions can be formulated for reactions between SiC and Mo that form the ternary compound. The overall reaction proceeds as Eq. (3) according to the phase diagram produced by Nowotny et al. This reaction can be broken down further into sub-reactions (4)–(6). 4:66Mo þ 2:65SiC/Mo4:66 Si2:65 C1:35 þ 1:30C
(3)
11Mo þ 3SiC/Mo5 Si3 þ 3Mo2 C
(4)
Mo5 Si3 þ 1:03Mo2 C þ 1:015SiC/1:515Mo4:66 Si2:65 C1:35
(5)
1:97Mo2 C þ 2:24SiC/0:845Mo4:66 Si2:65 C1:35 þ 3:07C
(6)
Eq. (4) is expected to proceed in the presence of high concentrations of molybdenum and will produce a mixture of molybdenum silicide and molybdenum carbide. The Mo–Si–C phase diagram by van Loo et al. [12] (1200 C isothermal section) indicates that Mo5Si3 can be in equilibrium with SiC at 1200 C. However, Mo2C is not stable with SiC at this temperature and is consumed through reaction (6). As identified in Fig. 1, although Mo5Si3 formed at temperatures of 1100 C and below, it reacted and disappeared by 1300 C. Reaction (5) is proposed for the consumption of Mo5Si3 and Mo2C to directly form the ternary compound. Analysis of powder samples from hot pressed Mo–SiC specimens revealed that only SiC and ternary compound were present (Fig. 2). The disparity in the relative intensities of several XRD peaks measured in this study for Mo4.66Si2.65C1.35 compared to those reported on the PDF card for Mo4.8Si3C0.6 is attributed to the different stoichiometries of the compounds. In this study, the ternary compound was in equilibrium with SiC and C, causing the C sites to be fully occupied as compared to Mo4.8Si3C0.6 from the PDF card, which is carbon deficient. Two polytypes of SiC, 3C and 6H, were identified in the hot pressed materials. The 3C polytype is a low temperature (<2000 C) phase, however, it was found to be the major SiC polytype for all specimens processed at 2000 and 2150 C. The combination of the nitrogen atmosphere and a liquid phase has been shown to stabilize or cause transformation from hexagonal polytypes into 3C SiC (nearly 100% 3C) at processing temperatures of 2300 C or higher [16–18]. Therefore, it appears that the processing conditions used in the present study were conducive to the stabilization of the 3C polytype.
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liquid formation in the SiC–Mo vertical section to be between 1900 and 1950 C [8]. 3.3. Thermal properties
Fig. 2. Powder X-ray diffraction representative of all hot pressed MS specimens.
3.2. Microstructure Geometric densities of the hot pressed specimens were 3.88, 3.88 and 4.08 g/cm3 for 50MS2000, 40MS2150 and 50MS2150, respectively. After reaction, the phase compositions were calculated to be (in vol%) 2.4% C, 23.8% Mo4.66Si2.65C1.35 and 73.8% SiC based on the stoichiometry of reaction (3) and an initial composition of 15 vol% Mo added to SiC. The theoretical density of the final composite was calculated to be 4.24 g/cm3 from the predicted composition and assuming theoretical densities of rC ¼ 2.22 g/cm3, rMo4:66 Si2:65 C1:35 ¼ 7:624 g=cm3 and rSiC ¼ 3.21 g/cm3. Analysis by SEM revealed the presence of three phases, SiC (grey), C (dark grey/black) and the ternary compound (light grey) (Fig. 3). Image analysis determined the volume percentages of the phases to be w73% SiC, w25% Mo4.66Si2.65C1.35 and w2% C for 50MS2150, which are in good agreement with the composition predicted from the reaction stoichiometry. The 50MS2000 had w73 vol% SiC as determined by image analysis. Determination of the ternary and carbon percentages was inhibited by the form the carbon was present. In materials processed above 2150 C, carbon was found as agglomerated platelets that were present within the ternary phase. For materials processed at 2000 C, carbon was present in contact with the ternary compound as flat sheets. However, the carbon did not form as well defined platelets as observed in MS2150 (both 40 and 50MS2150). Another difference to note is the large difference in the size of SiC particles. The microstructure of 50MS2000 is finer than that of the MS2150 materials, indicating an increase in mass transport leading to grain growth in the MS2150 specimens. This increased transport at elevated temperatures was consistent with an increase in the amount of liquid phase present during the final 1-h hold. Nowotny reported the first
Electrical resistivity was measured as a function of temperature to determine the impact of the electronic contribution on the thermal conductivity of the composites. Resistivity decreased from 4.62 U cm at room temperature to 4.15 U cm at 700 C. The resistivity is approximately six orders of magnitude higher than typical metals (Cu is w1.57 mU cm) [19], which makes the electronic contribution to thermal conductivity negligible based on the Weidemann–Franz law. Heat capacity of the MS2150 materials was measured from room temperature to 500 C and compared to the baseline SiC. The plot also contains the heat capacity calculated from the measured value for SiC and data reported [9] for Mo5Si2.22C1.16 (Fig. 4), as heat capacity has not been measured for Mo4.66Si2.65C1.35. The measured heat capacity of MS2150 ceramics increased from 0.54 J/g K at room temperature to 0.85 J/g K at 500 C. Heat capacity should only depend on the ratio of phases present (not the microstructure or distribution of phases), thus, the measured values for MS2150 were taken to be representative of all the MS materials. On average, the heat capacity of the MS material varied by only 2.5 1.7% from the predicted value over the temperature range of 50–500 C. Thermal diffusivities of both the MS2150 specimens were higher than the values measured for either the baseline SiC or Mo5Si3C1 (Fig. 5). The thermal diffusivities for the 50MS2000 and baseline SiC samples had similar values, varying from 2.1 to 5.5% from room temperature to 500 C. Thermal diffusivity values for 40MS2150 and 50MS2150 decreased from 50.5 and 46.8 mm2/s at room temperature to 11.6 and 11.3 mm2/s at 1000 C, respectively. This behavior is expected for composite materials in which the dominant heat transfer mechanism is phonon transport. In contrast, the diffusivity of Mo5Si3C1 increased with increasing temperature because the electron contribution dominated its thermal diffusivity [9]. The differences in the thermal diffusivity between the 50MS2000 and 40MS2150 samples were attributed to SiC grain size/interconnectivity and morphology of the graphite present. Thermal conductivities were calculated for all MS specimens and the baseline SiC from the measured thermal diffusivities, heat capacities, and bulk densities to 500 C without thermal expansion adjustment of the bulk density (Fig. 6). Two different MS2150 specimens had thermal conductivities in excess of 100 W/m K. The thermal conductivity of 50MS2150 was higher than the baseline SiC and 40MS2150 materials at all temperatures above 75 C due to its higher density. Thermal conductivity of 50MS2150 had a maximum value of 106 W/m K at 75 C and a minimum of w68 W/m K at
Fig. 3. Polished cross-sections showing the microstructures of 50MS2000 (left) and 50MS2150 (right) with the C, SiC and ternary phases labeled.
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1.30
1.00
Thermal Conductivity (W/m K)
1.10
Cp (J/g°C)
110
SiC (1 wt%Al) Mo5Si2.22C1.16 MS2150 Mixture
1.20
0.90 0.80 0.70 0.60 0.50 0.40 0.30
0
100
200
300
400
100 90 80 70 60 50
500
50MS2000 40MS2150 50MS2150 SiC (1 wt% Al)
Temperature (°C)
0
100
200
300
400
500
Temperature (°C)
Fig. 4. Heat capacity of MS2150 as compared to an SiC (1 wt% Al), Mo5Si2.22C1.16 [9], and the heat capacity predicted from measured heat capacities for SiC and Mo5Si2.22C1.16.
500 C. The 40MS2150 had a higher thermal conductivity than the baseline SiC below 175 C and had a maximum conductivity of 107 W/m K at room temperature. Thermal conductivity was approximately the same as the baseline SiC up to w400 C where the SiC seemed to reach a plateau value. The minimum thermal conductivity was 62 W/m K at 500 C. The thermal conductivity of 50MS2000 was the lowest of all the composite materials with a maximum of 92 W/m K at 75 C, decreasing to 56 W/m K at 500 C. The differences in diffusivity values between 50MS2000 and 40MS2150, which were attributed to differences in SiC grain size, resulted in the differences in thermal conductivity. Increases in thermal conductivity of the materials hot pressed with a liquid phase present are due to increased diffusivities, which were attributed to the increased SiC connectivity (and possibly increased grain size) in the MS2150 specimens. The increased diffusivity of the composite compared to that predicted based on the diffusivities of the constituent phases was attributed to the ability of the ternary compound to accommodate impurities, which removed impurities that could act as phonon scattering sites from SiC grains or grain boundaries. Eucken’s model [21] was used to analyze the thermal conductivities of the MS ceramics to verify that the enhanced thermal conductivities of the composite were due to increased conductivity of the SiC. A rearrangement of the Eucken model is presented as Eq. (7) and assumes a continuous phase and a dispersed phase where
Fig. 6. Thermal conductivity of Mo–SiC ceramics with that of an SiC (1 wt% Al).
km ¼ kc
1 þ 2Vd ð1 kc =kd Þ=ð2kc =kd þ 1Þ 1 Vd ð1 kc =kd Þ=ðkc =kd þ 1Þ
(7)
where km is the thermal conductivity of the mixture, kc is the thermal conductivity of the continuous phase (SiC), kd is the thermal conductivity of the dispersed phase and Vd is the volume fraction of dispersed phase. Eq. (7) allows calculation of the thermal conductivity of the SiC in the MS materials. The results of these calculations revealed that the continuous SiC in the MS materials had a room temperature thermal conductivity between 142 and 167 W/m K. For comparison, the room temperature thermal conductivity calculated for the baseline SiC material was 105 W/m K. For 40MS2150 (Fig. 7), the increase in the thermal conductivity of the SiC as compared to the baseline material was w60% at room temperature and w30% at 500 C. The average increase in thermal conductivity from room temperature to 500 C was w30% for 50MS2000 and w58% for 50MS2150. Because the ternary compound has a lower thermal conductivity than SiC, the initial expectation was that its addition to SiC would decrease the overall thermal conductivity of the composite material. However, the ternary compound has the ability to accommodate oxygen and nitrogen into its structure [20]. One explanation for the increased thermal diffusivities and conductivities of the Mo–SiC ceramics over the baseline SiC ceramic is that impurities could preferentially segregate from the SiC and the grain boundaries to the liquid present during processing and become
180
Thermal Conductivity (W/m K)
Thermal Diffusivity (mm2/s)
60 50MS2000 40MS2150 50MS2150 SiC (1 wt% Al) Mo5Si2.22C1.16
50 40 30 20 10 0
160 140 120 100 80
40 20 0
0
100
200
300
400
500
600
700
800
900
1000
Mo5Si2.22C1.16 SiC (1 wt% Al) SiC 40MS2150 40MS2150
60
0
100
200
300
400
500
Temperature (°C)
Temperature (°C) Fig. 5. Thermal diffusivity of MS ceramics as compared to SiC (1 wt% Al) and Mo5 Si2.22C1.16 [9].
Fig. 7. Prediction of the thermal conductivity of SiC in 40MS2150 using the Eucken model compared to the calculated thermal conductivities of 40MS2150, the baseline SiC composition, and Mo5Si2.22C1.16 [9].
A.A. Buchheit et al. / Intermetallics 16 (2008) 854–859
trapped in the ternary phase upon solidification thereby reducing phonon scattering in the higher conductivity phase (SiC) and increasing its thermal conductivity. Verification of this hypothesis would require more detailed analysis (e.g., transmission electron microscopy), which was not performed in the current study. 4. Conclusions Composite ceramics of SiC and an Mo–Si–C ternary compound were produced by reaction hot pressing. The composites reached up to 96% of their theoretical density. Residual carbon from the reaction between Mo and SiC was found to be present as inclusions in the ternary compound in specimens processed at 2150 C. Phase evolution of the Mo–SiC reaction was examined from 900 to 1900 C. A finding of note was the stability of Mo5Si3 with SiC found by van Loo et al. at 1200 C extended up to 1300 C. The ternary Mo–Si–C compound was shown to be a promising additive for the fabrication of SiC-based ceramics by hot pressing reaction at 2150 C. Despite the low thermal conductivity of the ternary phase (w11 W/m K at room temperature), the resulting ceramics had thermal conductivities equal to or higher than a baseline SiC composition (92, 107 and 104 W/m K at room temperature for 50MS2000, 40MS2150 and 50MS2150, respectively). Reducing the amount of the ternary phase may lead to further increases in the thermal conductivity of the composite, while still allowing for liquid phase densification. Acknowledgements This work was supported by the Army Space and Missile Defense Command under grant number BAA DASG60-00-0005 and in part by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Transportation Technologies as part of the High Temperature Materials Laboratory User Program at Oak Ridge National Laboratory managed by the UT-Battelle LLC, for the Department of Energy under contract DE-AC05000OR22725. Financial support for A. Buchheit was provided by the Department of Education under the Graduate Assistance in Areas of National Need (GAANN) Fellowship. The authors would also like to thank Ms. Jennifer Gilmore for technical assistance.
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References [1] Riedel R, editor. Handbook of ceramic hard materials, vol. 2. Wiley-VCH Verlag GmbH; 2000. p. 721. [2] Zhu Q, Shobu K. SiC–Mo5Si3C1 composites by melt infiltration process. J Mater Sci Lett 2000;19:153–5. [3] Hayashi T, Ito K, Tanaka K. Physical and mechanical properties of single crystals of the Mo5Si3C phase. Intermetallics 2003;11:835–40. [4] Peng LM, Wang JH, Li J, He LH, Zhao JH. Synthesis and mechanical properties of ternary molybdenum carbosilicide and its composite. J Mater Sci Lett 2005;40: 2705–7. [5] Zhu Q, Shobu K, Tani E, Kishi K, Umebayashi S. High-temperature strength and creep behavior of melt-infiltrated SiC–Mo5Si3C1 composites. J Am Ceram Soc 1999;82(8):2276–8. [6] Zhu Q, Shobu K. High-temperature fracture toughness of SiC–Mo5Si3C composites. J Mater Sci Lett 2000;19:1529–31. [7] Zhu Q, Shobu K. High-temperature mechanical properties of SiC–Mo5(Si,Al)3C composites. J Am Ceram Soc 2001;84(2):413–9. [8] Nowotny H, Parthe E, Kieffer R, Benesovsky F. The three-component system Mo–Si–C. Monatsh Chem 1954;85:255–72. [9] Buchheit A, Fahrenholtz W, Hilmas G, Deason D, Wang H. Thermal properties of a hot pressed Mo5Si3C1 ceramic, Intermetallics, in press. [10] North B, Gilchrist K. Effect of impurity doping on a reaction-bonded silicon carbide. Ceram Bull 1981;60(5):549–54. [11] Keppler U. Thermal expansion of silicon. Z Metallkd 1988;79(3):157–8. [12] van Loo F, Smet F, Rieck G, Gerard D, Verspui G. Phase relations and diffusion paths in the molybdenum–silicon–carbon system at 1200 C. High Temp. – High Pressures 1982;14(1):25–31. [13] Martinelli A, Drew R, Berriche R. Correlation between the strength of SiC–Mo diffusion couples and the mechanical properties of the interfacial reaction products. J Mater Sci Lett 1996;15:307–10. [14] Morozumi S, Kikuchi M, Sugai S, Hayahsi M. Reaction between molybdenum and carbon, and several carbides. Nippon Kinzoku Gakkaishi 1980;44(12): 1404–13. [15] Suzuki Y, Niihara K. Synthesis and mechanical properties of Mo5Si3C1 and Mo5Si3C1-based composites. Intermetallics 1998;6:7–13. [16] Turan S, Knowles KM. a / b Reverse phase transformation in silicon carbide in silicon nitride–particulate-reinforced-silicon carbide composites. J Am Ceram Soc 1996;79(11):2892–6. [17] Kieffer AR, Ettmayer P, Gugel E, Schmidt A. In: Henisch HR, Roy R, editors. Silicon carbide-1968. Elmsford, NY: Pergamon; 1969. p. S153–66. [18] Turan S, Knowles KM. A comparison of the microstructure of silicon nitride– silicon carbide composites made with and without deoxidized starting material. J Microsc 1995;177(3):287–304. [19] Kasap SO. Principles of electronic materials and devices. 3rd ed. Boston: McGraw Hill; 2006. [20] Costa e Silva A, Kaufman M. Phase relations in the Mo–Si–C system relevant to the processing of MoSi2–SiC composites. Metall Mater Trans A 1994;25A(1): 5–15. [21] Kingery WD, Bowen HK, Uhlmann DR. Introduction to ceramics. New York: John Wiley & Sons; 1976.
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Processing and thermal properties of an Mo5Si3C–SiC ceramic Andrew A. Buchheit a, Greg E. Hilmas b, *, William G. Fahrenholtz b, Douglas M. Deason c, Hsin Wang d a
Coorstek, Inc., Golden, CO 80401, USA Materials Science and Engineering Department, University of Missouri-Rolla, Rolla, MO 65409, USA c U.S. Army Space & Missile Defense Command, Building 5220, Redstone Arsenal, AL 35898, USA d Oak Ridge National Laboratory, High Temperature Materials Laboratory, Oak Ridge, TN 37831, USA b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 10 July 2007 Received in revised form 25 February 2008 Accepted 8 March 2008 Available online 3 June 2008
Composite ceramics containing w25 vol% of the Mo–Si–C ternary compound and SiC (referred to as MS) were reaction hot pressed up to 96% relative density. The electrical resistivity of the composite processed at 2150 C decreased from w4.60 U cm at room temperature to 4.15 U cm at 700 C. Thermal conductivity of the MS ceramics processed at 2150 C was above 100 W/m K at room temperature, decreasing to between 62 and 68 W/m K at 500 C. Using the Eucken model for thermal conductivity, the interconnected SiC phase in the MS materials was calculated to have a room temperature thermal conductivity between 160 and 170 W/m K. As compared to a baseline SiC composition, the continuous SiC in the MS materials had an average increase in thermal conductivity of w50% over the temperature range of room temperature to 500 C. This increase in thermal conductivity was attributed to the accommodation of impurities that would typically be present in SiC grains and grain boundaries (e.g., N and O) into the ternary phase. Ó 2008 Elsevier Ltd. All rights reserved.
Keywords: A. Molybdenum silicides B. Thermal properties C. Reaction synthesis F. Diffraction F. Electron microscopy, scanning
1. Introduction
is written as Mo5Si3C and crystallizes in the D88 structure (P63/ Stoichiometry of the ternary compound ranges from 50 to 60.7% molybdenum, 27.1 to 36.8% silicon and 8.3 to 15.6% carbon based on the Nowotny phase diagram and the eutectic temperature in the SiC–C-ternary phase field was reported to be w1900 C [8]. The composition Mo4.89Si3C0.686, which is within the ternary compound stoichiometry, has a coefficient of thermal expansion (a) anisotropy (ac-axis/aa-axis) of w2 with ac-axis having a room temperature value of 9.2 ppm/ C. The single crystal thermal conductivity of the ternary compound with the same stoichiometry was reported to be 8.5 W/m K in the ½1120 direction. A room temperature conductivity of w10 W/m K was measured for polycrystalline Mo5Si2.22C1.16, which is the composition stable in contact with C and MoC [9]. In contrast to the ternary compound, silicon has a thermal conductivity of w102 W/m K and an isotropic CTE of w4.4 ppm/ C from room temperature to 900 C [10,11]. As a result, Si melt infiltrated SiC has a thermal conductivity between 100 and 175 W/m K at room temperature [10]. With these differences in property values between the ternary Mo–Si–C phase and Si, the thermal conductivity of the SiC–Mo5Si3C ceramics is expected to be lower than that of Si-infiltrated SiC ceramics. The objective of this paper is to characterize the thermal properties of SiC–Mo5Si3C ceramics in addition to the mechanism of ternary phase formation. The in situ reaction and densification process, resulting in an overall composition in the range of those studied by Zhu et al., were explored as an alternative to melt infiltration. mcm).
The properties of silicon carbide ceramics depend on the processing method. Melt infiltration is one of the methods employed to manufacture SiC ceramics. Generally, the infiltrant is silicon, which limits the use of melt infiltrated materials to temperatures below the melting temperature of silicon (w1400 C). Average strengths of Si-infiltrated SiC are reported to be w350 MPa at room temperature and 200 MPa at 1400 C [1]. Within the past decade, Zhu et al. studied the ternary Mo–Si–C compound as an alternative infiltrant [2] due to its high melting temperature (w2100 C) and its ductile to brittle transition around 1300 C [3–5]. To date, mechanical properties have been measured for compositions containing 70 vol% SiC with 25–30 vol% of the ternary compound [5,6]. From the latter studies, approximately 90% of the room temperature strength (290 MPa for 92% dense materials) was retained up to 1700 C [7] and fracture toughness increased up to w8 MPa m1/2 above the ductile to brittle transition temperature of the ternary compound [6]. The properties of the ternary Mo–Si–C compound are significantly different than those of silicon. The stoichiometric compound
* Corresponding author. Materials Science and Engineering Department, University of Missouri-Rolla, 225 McNutt Hall, Rolla, MO 65401, USA. Tel.: þ1 573 341 7205. E-mail addresses: [email protected] (A.A. Buchheit), [email protected] (G.E. Hilmas). 0966-9795/$ – see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2008.03.010
A.A. Buchheit et al. / Intermetallics 16 (2008) 854–859
2. Experimental procedure 2.1. Powder processing and specimen fabrication Molybdenum powder was 99.95% pure and had an average particle size in the range of 2–4 mm (M-1109, Cerac, Inc., Milwaukee, WI). Silicon carbide powder was >95% b-SiC (BF-12, H.C. Starck, Goslar, Germany) with a reported average particle size of 0.6 mm and a surface area of 11.0–13.0 m2/g. The reactant mixture contained 15 vol% molybdenum with the remainder being SiC. Powders were attrition milled (Model 01-HD, Union Process, Akron, OH) for 1 h in hexane at 600 rpm using a fluoro-polymer coated spindle, 750 cc bucket and 3/16’’ (4.76 mm) diameter SiC media. Powders ¨ CHI Laborwere dried by rotary evaporation (Rotavapor R-124, BU technik AG, Flawil, Switzerland) at 70 C, 75 rpm and 40 Pa vacuum. Specimens will be referred to as MS with a numerical prefix designating the processing pressure (in MPa) and a numerical suffix designating the processing temperature (in C), as shown in Table 1. The batch has an overall composition that falls in a three phase field containing SiC, C, and the ternary compound. Several assumptions were made when predicting the stoichiometry of the ternary compound from the atomic percentages of Mo, Si and C in the batch. The first assumption was that molybdenum and silicon in the ternary compound only resided on Mo and Si lattice sites, and the maximum number of sites in the unit cell was 5 and 3, respectively. Carbon was assumed to occupy all of the C sites, but also substitute onto Si sites due to its similar bonding characteristics and smaller atomic size compared to Si. This gives the constraint that C þ Si must be 4. Following these rules and using the Nowotny Mo–Si–C phase diagram [8] the stoichiometry of the ternary compound in equilibrium with SiC and C was predicted to be 53.81% Mo, 30.60% Si and 15.59%, which corresponds to Mo4.66Si2.65C1.35. A graphite die was lined with graphite foil (UCAR Carbon Company, Inc., Columbia, TN). Powder was added to the die and it was pressed uniaxially at 15 MPa (Model C, Carver, Inc., Wabash, IN). Specimens were processed in a graphite element hot press (HP20-3060, Thermal Technology, Inc., Santa Rosa, CA). During the hot pressing cycle, vacuum was maintained at <16 Pa up to 1950 C and then the furnace was backfilled with helium. Above 1950 C, a flowing He atmosphere was maintained at a slightly positive pressure (w105 Pa). At 1100 C, a uniaxial pressure of 25 MPa was applied and the pressure was increased to either 40 or 50 MPa at 1900 C, where temperature was held for 30 min. The final hold was 1 h at either 2000 or 2150 C. Cooling occurred at w30 /min until 1500 C. The pressure was released when the furnace temperature reached 1750 C. Below 1500 C, the furnace was allowed to cool uncontrolled. Three different combinations of processing conditions were used to fabricate specimens (Table 1). Bulk densities were determined geometrically using disks w12.5 mm in diameter and w3 mm thick. For comparison, a baseline alpha SiC (UF-10, H.C. Starck, Goslar, Germany) composition containing 1 wt% Al and 3 wt% C was hot pressed at 2000 C and 32 MPa for 1 h. 2.2. XRD and SEM analyses Phase evolution was investigated by X-ray diffraction analysis (XRD; Scintag XDS 2000, Thermo Optek Corporation, Franklin,
Table 1 Specimen designations and the corresponding processing conditions Designation
Pressure (MPa)
Temperature ( C)
50MS2000 40MS2150 50MS2150 SiC (1 wt% Al)
50 40 50 32
2000 2150 2150 2000
855
MA) on powder samples that were heated in a graphite furnace (Model 1000-3060-FP20, Thermal Technology, Inc., Santa Rosa, CA) with the same temperature and atmosphere profile as those hot pressed, except that temperature did not exceed 1900 C. The final composition of the hot pressed (HP) specimens was determined using XRD analysis of powders produced by crushing and grinding hot pressed materials. Powders were scanned from 20 to 80 2q using Cu Ka radiation. Patterns were indexed to powder diffraction file cards for 3C SiC (29-1129), 6H SiC (29-1131), Mo2C (35-0787), Mo5Si3 (34-0371) and Mo4.8Si3C0.6 (43-1199). Specimens were sectioned perpendicular to the hot pressing direction and polished to 3 mm for examination by scanning electron microscopy (SEM; S-570 SEM, Hitachi, Tokyo, Japan). Multiple images of polished surfaces were examined using image analysis software (ImageJ, National Institute of Health, Bethesda, MD) to estimate volume percentages of phases present in the hot pressed specimens.
2.3. Thermal testing Electrical resistance was measured using a four point probe potientiodynamic setup (Solartron 1470, Solartron Analytical, Farnborough, Hampshire, England). Bars w20 mm 1.5 mm 2 mm were notched in four places around their perimeters to allow attachment of platinum electrode wires. The wires were wrapped around the specimens and silver paste was applied to improve contact. Specimens were heated to 700 C under flowing argon to sinter the silver paste. During cooling, resistance measurements were recorded every 50 C between 700 and 400 C and then every 100 C to room temperature. Electrical resistivity was calculated using the following equation:
r ¼
RA L
(1)
where r is resistivity, R is the measured resistance, A is the crosssectional area of the specimen and L is the gauge length between the two voltage (inner) leads. Heat capacity was measured every 50 C between room temperature and 500 C in flowing high-purity argon with a ramp rate of 20 C/min (STA 409C, Netzsch, Selb/Bavaria Germany). A sapphire disk was used as a standard and all measurements were taken during the heating portion of the cycle. Data reported are from specimens w5 mm in diameter and <1 mm thickness. The thermal diffusivities of 50MS2000, 40MS2150 and the baseline SiC were measured at Oak Ridge National Laboratory (ORNL) in the High Temperature Materials Laboratory (HTML) user facility using a commercial laser flash system (Flashline 5000, Anter Corp., Pittsburgh, PA) on w12.5 mm diameter disks of w3 mm thickness. The reported thermal diffusivities were the average of three measurements made every w100 C and based on the Clark and Taylor data (ASTM E1461-01). Two separate furnaces with temperature ranges of RT to 500 C and 500 to >1900 C were used to determine the diffusivity to 1000 C. A silicon-based infrared pyrometer was used for temperature measurements. The thermal diffusivities of specimens of compositions 40MS2150 and 50MS2150, the same size as those tested at ORNL, were also measured every 100 C from room temperature to 900 C (LFA 457, Netzsch, Selb/Bavaria Germany) using a system equipped with an InSb detector. Thermal conductivities were then calculated using Eq. (2) where d is the thermal diffusivity, r is the bulk density and CP is the heat capacity. k ¼ drCP
(2)
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3. Results and discussion 3.1. Phase evolution The phase evolution study revealed molybdenum and SiC did not react below 700 C (Fig. 1). At 900 and 1100 C, Mo5Si3, Mo2C and a trace amount of Mo5Si3C1 (at 1100 C only) were detected along with SiC. A study by van Loo et al. [12] on the Mo–Si–C system reported the formation of Mo5Si3, Mo2C and Mo5Si3C at 1200 C, which agrees well with the analysis at 1100 C in the present study. At 1300 C, both Mo5Si3C1 and Mo2C were detected. Above 1750 C, only SiC and Mo5Si3C1 were present. Literature regarding SiC–Mo diffusion couples found the formation of both Mo2C and Mo5Si3 by 1200 C and that Mo5Si3 converted to Mo5Si3C1 between 1400 and 1500 C [13]. A previous study of SiC–Mo diffusion couples reported that only Mo2C and Mo3Si2 (Mo5Si3) formed when the couple was held for 9 h at 1700 C [14]. Based on the electron microprobe analysis performed by Morozumi et al. and the Nowotny Mo–Si–C phase diagram [8], that specimen had a layered structure with, most likely, a layer of Mo5Si3C1 in contact with SiC, a layer containing a mixture of Mo5Si3 and Mo2C, and then a layer of Mo2C or MoC1x in contact with Mo. The phase evolution study revealed that between the temperature ranges of the van Loo and Nowotny phase diagrams (1200–1600 C) Mo5Si3 and SiC were no longer in equilibrium. This was shown to occur by 1300 C with only a 30-min hold in the present study as no Mo5Si3 was detected by XRD. Also, above 1750 C, no detectable amount of Mo2C was found. Rietveld analysis of the XRD patterns, collected as part of the present study for powders heated to temperatures from 1300 to 1900 C (Fig. 1), was used to determine the lattice parameters of the ternary compound with a stoichiometry of Mo4.66Si2.65C1.35 (Table 2). Different stoichiometries of the ternary phase have been examined and lattice parameters have been reported in the literature [15]. Reported lattice parameters for Mo4.7Si3C0.6 hot pressed at 1500 C are 7.290 Å for the a and b axes and 5.043 Å for the c-axis [15]. The lattice parameters have also been reported for the ternary compound without specific stoichiometries, although the ternary compound was in equilibrium with either Mo2C (ranges from Mo5Si2.19C1.05 to Mo5Si2.66C0.69) or MoSi2 (ranges from Mo3.93Si2.96C1.04 to Mo4.31Si3C0.69), which will be referred to as Mo-rich and Mo-poor regions, respectively. The a-axis lattice parameter for the Mo-rich ternary compound was 7.2878 Å and the c-axis parameter was 5.5069 Å. For the Mo-poor compound, the a-axis parameter was 7.2884 Å and the c-axis parameter was 5.0164 Å [15]. The c-axis parameter calculated in the present study falls within previous ranges reported in the literature. Notice the
Fig. 1. Phase evolution in Mo–SiC mixtures between 700 and 1900 C with selected phase-specific peaks labeled.
Table 2 Lattice parameters and crystal volume of Mo4.66Si2.65C1.35 from the phase evolution study at their corresponding temperatures Temperature
a/b axis (Å)
c-axis (Å)
c/a
Volume (Å3)
1300 1500 1750 1900
7.26(4) 7.26(8) 7.26(7) 7.26(8)
5.04(7) 5.04(4) 5.03(8) 5.03(9)
0.695 0.694 0.693 0.693
230.67 230.73 230.43 230.60
Mo stoichiometry is nearly the same as that reported by Suzuki et al. (Mo4.66 vs. Mo4.7) and the average measured c-axis was 5.042 0.004 Å, suggesting the cell height was controlled by the Mo–Si bonds. However, the lower a-axis parameter was attributed to the substitution of C onto Si sites in combination with fully occupied C sites, which should result in a decrease in the size of the Mo–Si octahedra. The lattice parameter analysis was performed to allow the theoretical density to be calculated. Using the lattice parameters, the crystal volume of Mo4.66Si2.65C1.35 was calculated to be 230.61 0.13 Å3 with a molar weight of 529.58 g/ mol. These values result in a theoretical density of 7.624 g/cm3 for Mo4.66Si2.65C1.35 as compared to 7.89 g/cm3 as reported in PDF card 43-1199 for Mo4.8Si3C0.6. The decrease in Mo content and the increase in C content of the material examined in this study led to a slightly lower theoretical density despite the decrease in unit cell volume. Several chemical expressions can be formulated for reactions between SiC and Mo that form the ternary compound. The overall reaction proceeds as Eq. (3) according to the phase diagram produced by Nowotny et al. This reaction can be broken down further into sub-reactions (4)–(6). 4:66Mo þ 2:65SiC/Mo4:66 Si2:65 C1:35 þ 1:30C
(3)
11Mo þ 3SiC/Mo5 Si3 þ 3Mo2 C
(4)
Mo5 Si3 þ 1:03Mo2 C þ 1:015SiC/1:515Mo4:66 Si2:65 C1:35
(5)
1:97Mo2 C þ 2:24SiC/0:845Mo4:66 Si2:65 C1:35 þ 3:07C
(6)
Eq. (4) is expected to proceed in the presence of high concentrations of molybdenum and will produce a mixture of molybdenum silicide and molybdenum carbide. The Mo–Si–C phase diagram by van Loo et al. [12] (1200 C isothermal section) indicates that Mo5Si3 can be in equilibrium with SiC at 1200 C. However, Mo2C is not stable with SiC at this temperature and is consumed through reaction (6). As identified in Fig. 1, although Mo5Si3 formed at temperatures of 1100 C and below, it reacted and disappeared by 1300 C. Reaction (5) is proposed for the consumption of Mo5Si3 and Mo2C to directly form the ternary compound. Analysis of powder samples from hot pressed Mo–SiC specimens revealed that only SiC and ternary compound were present (Fig. 2). The disparity in the relative intensities of several XRD peaks measured in this study for Mo4.66Si2.65C1.35 compared to those reported on the PDF card for Mo4.8Si3C0.6 is attributed to the different stoichiometries of the compounds. In this study, the ternary compound was in equilibrium with SiC and C, causing the C sites to be fully occupied as compared to Mo4.8Si3C0.6 from the PDF card, which is carbon deficient. Two polytypes of SiC, 3C and 6H, were identified in the hot pressed materials. The 3C polytype is a low temperature (<2000 C) phase, however, it was found to be the major SiC polytype for all specimens processed at 2000 and 2150 C. The combination of the nitrogen atmosphere and a liquid phase has been shown to stabilize or cause transformation from hexagonal polytypes into 3C SiC (nearly 100% 3C) at processing temperatures of 2300 C or higher [16–18]. Therefore, it appears that the processing conditions used in the present study were conducive to the stabilization of the 3C polytype.
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liquid formation in the SiC–Mo vertical section to be between 1900 and 1950 C [8]. 3.3. Thermal properties
Fig. 2. Powder X-ray diffraction representative of all hot pressed MS specimens.
3.2. Microstructure Geometric densities of the hot pressed specimens were 3.88, 3.88 and 4.08 g/cm3 for 50MS2000, 40MS2150 and 50MS2150, respectively. After reaction, the phase compositions were calculated to be (in vol%) 2.4% C, 23.8% Mo4.66Si2.65C1.35 and 73.8% SiC based on the stoichiometry of reaction (3) and an initial composition of 15 vol% Mo added to SiC. The theoretical density of the final composite was calculated to be 4.24 g/cm3 from the predicted composition and assuming theoretical densities of rC ¼ 2.22 g/cm3, rMo4:66 Si2:65 C1:35 ¼ 7:624 g=cm3 and rSiC ¼ 3.21 g/cm3. Analysis by SEM revealed the presence of three phases, SiC (grey), C (dark grey/black) and the ternary compound (light grey) (Fig. 3). Image analysis determined the volume percentages of the phases to be w73% SiC, w25% Mo4.66Si2.65C1.35 and w2% C for 50MS2150, which are in good agreement with the composition predicted from the reaction stoichiometry. The 50MS2000 had w73 vol% SiC as determined by image analysis. Determination of the ternary and carbon percentages was inhibited by the form the carbon was present. In materials processed above 2150 C, carbon was found as agglomerated platelets that were present within the ternary phase. For materials processed at 2000 C, carbon was present in contact with the ternary compound as flat sheets. However, the carbon did not form as well defined platelets as observed in MS2150 (both 40 and 50MS2150). Another difference to note is the large difference in the size of SiC particles. The microstructure of 50MS2000 is finer than that of the MS2150 materials, indicating an increase in mass transport leading to grain growth in the MS2150 specimens. This increased transport at elevated temperatures was consistent with an increase in the amount of liquid phase present during the final 1-h hold. Nowotny reported the first
Electrical resistivity was measured as a function of temperature to determine the impact of the electronic contribution on the thermal conductivity of the composites. Resistivity decreased from 4.62 U cm at room temperature to 4.15 U cm at 700 C. The resistivity is approximately six orders of magnitude higher than typical metals (Cu is w1.57 mU cm) [19], which makes the electronic contribution to thermal conductivity negligible based on the Weidemann–Franz law. Heat capacity of the MS2150 materials was measured from room temperature to 500 C and compared to the baseline SiC. The plot also contains the heat capacity calculated from the measured value for SiC and data reported [9] for Mo5Si2.22C1.16 (Fig. 4), as heat capacity has not been measured for Mo4.66Si2.65C1.35. The measured heat capacity of MS2150 ceramics increased from 0.54 J/g K at room temperature to 0.85 J/g K at 500 C. Heat capacity should only depend on the ratio of phases present (not the microstructure or distribution of phases), thus, the measured values for MS2150 were taken to be representative of all the MS materials. On average, the heat capacity of the MS material varied by only 2.5 1.7% from the predicted value over the temperature range of 50–500 C. Thermal diffusivities of both the MS2150 specimens were higher than the values measured for either the baseline SiC or Mo5Si3C1 (Fig. 5). The thermal diffusivities for the 50MS2000 and baseline SiC samples had similar values, varying from 2.1 to 5.5% from room temperature to 500 C. Thermal diffusivity values for 40MS2150 and 50MS2150 decreased from 50.5 and 46.8 mm2/s at room temperature to 11.6 and 11.3 mm2/s at 1000 C, respectively. This behavior is expected for composite materials in which the dominant heat transfer mechanism is phonon transport. In contrast, the diffusivity of Mo5Si3C1 increased with increasing temperature because the electron contribution dominated its thermal diffusivity [9]. The differences in the thermal diffusivity between the 50MS2000 and 40MS2150 samples were attributed to SiC grain size/interconnectivity and morphology of the graphite present. Thermal conductivities were calculated for all MS specimens and the baseline SiC from the measured thermal diffusivities, heat capacities, and bulk densities to 500 C without thermal expansion adjustment of the bulk density (Fig. 6). Two different MS2150 specimens had thermal conductivities in excess of 100 W/m K. The thermal conductivity of 50MS2150 was higher than the baseline SiC and 40MS2150 materials at all temperatures above 75 C due to its higher density. Thermal conductivity of 50MS2150 had a maximum value of 106 W/m K at 75 C and a minimum of w68 W/m K at
Fig. 3. Polished cross-sections showing the microstructures of 50MS2000 (left) and 50MS2150 (right) with the C, SiC and ternary phases labeled.
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1.30
1.00
Thermal Conductivity (W/m K)
1.10
Cp (J/g°C)
110
SiC (1 wt%Al) Mo5Si2.22C1.16 MS2150 Mixture
1.20
0.90 0.80 0.70 0.60 0.50 0.40 0.30
0
100
200
300
400
100 90 80 70 60 50
500
50MS2000 40MS2150 50MS2150 SiC (1 wt% Al)
Temperature (°C)
0
100
200
300
400
500
Temperature (°C)
Fig. 4. Heat capacity of MS2150 as compared to an SiC (1 wt% Al), Mo5Si2.22C1.16 [9], and the heat capacity predicted from measured heat capacities for SiC and Mo5Si2.22C1.16.
500 C. The 40MS2150 had a higher thermal conductivity than the baseline SiC below 175 C and had a maximum conductivity of 107 W/m K at room temperature. Thermal conductivity was approximately the same as the baseline SiC up to w400 C where the SiC seemed to reach a plateau value. The minimum thermal conductivity was 62 W/m K at 500 C. The thermal conductivity of 50MS2000 was the lowest of all the composite materials with a maximum of 92 W/m K at 75 C, decreasing to 56 W/m K at 500 C. The differences in diffusivity values between 50MS2000 and 40MS2150, which were attributed to differences in SiC grain size, resulted in the differences in thermal conductivity. Increases in thermal conductivity of the materials hot pressed with a liquid phase present are due to increased diffusivities, which were attributed to the increased SiC connectivity (and possibly increased grain size) in the MS2150 specimens. The increased diffusivity of the composite compared to that predicted based on the diffusivities of the constituent phases was attributed to the ability of the ternary compound to accommodate impurities, which removed impurities that could act as phonon scattering sites from SiC grains or grain boundaries. Eucken’s model [21] was used to analyze the thermal conductivities of the MS ceramics to verify that the enhanced thermal conductivities of the composite were due to increased conductivity of the SiC. A rearrangement of the Eucken model is presented as Eq. (7) and assumes a continuous phase and a dispersed phase where
Fig. 6. Thermal conductivity of Mo–SiC ceramics with that of an SiC (1 wt% Al).
km ¼ kc
1 þ 2Vd ð1 kc =kd Þ=ð2kc =kd þ 1Þ 1 Vd ð1 kc =kd Þ=ðkc =kd þ 1Þ
(7)
where km is the thermal conductivity of the mixture, kc is the thermal conductivity of the continuous phase (SiC), kd is the thermal conductivity of the dispersed phase and Vd is the volume fraction of dispersed phase. Eq. (7) allows calculation of the thermal conductivity of the SiC in the MS materials. The results of these calculations revealed that the continuous SiC in the MS materials had a room temperature thermal conductivity between 142 and 167 W/m K. For comparison, the room temperature thermal conductivity calculated for the baseline SiC material was 105 W/m K. For 40MS2150 (Fig. 7), the increase in the thermal conductivity of the SiC as compared to the baseline material was w60% at room temperature and w30% at 500 C. The average increase in thermal conductivity from room temperature to 500 C was w30% for 50MS2000 and w58% for 50MS2150. Because the ternary compound has a lower thermal conductivity than SiC, the initial expectation was that its addition to SiC would decrease the overall thermal conductivity of the composite material. However, the ternary compound has the ability to accommodate oxygen and nitrogen into its structure [20]. One explanation for the increased thermal diffusivities and conductivities of the Mo–SiC ceramics over the baseline SiC ceramic is that impurities could preferentially segregate from the SiC and the grain boundaries to the liquid present during processing and become
180
Thermal Conductivity (W/m K)
Thermal Diffusivity (mm2/s)
60 50MS2000 40MS2150 50MS2150 SiC (1 wt% Al) Mo5Si2.22C1.16
50 40 30 20 10 0
160 140 120 100 80
40 20 0
0
100
200
300
400
500
600
700
800
900
1000
Mo5Si2.22C1.16 SiC (1 wt% Al) SiC 40MS2150 40MS2150
60
0
100
200
300
400
500
Temperature (°C)
Temperature (°C) Fig. 5. Thermal diffusivity of MS ceramics as compared to SiC (1 wt% Al) and Mo5 Si2.22C1.16 [9].
Fig. 7. Prediction of the thermal conductivity of SiC in 40MS2150 using the Eucken model compared to the calculated thermal conductivities of 40MS2150, the baseline SiC composition, and Mo5Si2.22C1.16 [9].
A.A. Buchheit et al. / Intermetallics 16 (2008) 854–859
trapped in the ternary phase upon solidification thereby reducing phonon scattering in the higher conductivity phase (SiC) and increasing its thermal conductivity. Verification of this hypothesis would require more detailed analysis (e.g., transmission electron microscopy), which was not performed in the current study. 4. Conclusions Composite ceramics of SiC and an Mo–Si–C ternary compound were produced by reaction hot pressing. The composites reached up to 96% of their theoretical density. Residual carbon from the reaction between Mo and SiC was found to be present as inclusions in the ternary compound in specimens processed at 2150 C. Phase evolution of the Mo–SiC reaction was examined from 900 to 1900 C. A finding of note was the stability of Mo5Si3 with SiC found by van Loo et al. at 1200 C extended up to 1300 C. The ternary Mo–Si–C compound was shown to be a promising additive for the fabrication of SiC-based ceramics by hot pressing reaction at 2150 C. Despite the low thermal conductivity of the ternary phase (w11 W/m K at room temperature), the resulting ceramics had thermal conductivities equal to or higher than a baseline SiC composition (92, 107 and 104 W/m K at room temperature for 50MS2000, 40MS2150 and 50MS2150, respectively). Reducing the amount of the ternary phase may lead to further increases in the thermal conductivity of the composite, while still allowing for liquid phase densification. Acknowledgements This work was supported by the Army Space and Missile Defense Command under grant number BAA DASG60-00-0005 and in part by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Transportation Technologies as part of the High Temperature Materials Laboratory User Program at Oak Ridge National Laboratory managed by the UT-Battelle LLC, for the Department of Energy under contract DE-AC05000OR22725. Financial support for A. Buchheit was provided by the Department of Education under the Graduate Assistance in Areas of National Need (GAANN) Fellowship. The authors would also like to thank Ms. Jennifer Gilmore for technical assistance.
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References [1] Riedel R, editor. Handbook of ceramic hard materials, vol. 2. Wiley-VCH Verlag GmbH; 2000. p. 721. [2] Zhu Q, Shobu K. SiC–Mo5Si3C1 composites by melt infiltration process. J Mater Sci Lett 2000;19:153–5. [3] Hayashi T, Ito K, Tanaka K. Physical and mechanical properties of single crystals of the Mo5Si3C phase. Intermetallics 2003;11:835–40. [4] Peng LM, Wang JH, Li J, He LH, Zhao JH. Synthesis and mechanical properties of ternary molybdenum carbosilicide and its composite. J Mater Sci Lett 2005;40: 2705–7. [5] Zhu Q, Shobu K, Tani E, Kishi K, Umebayashi S. High-temperature strength and creep behavior of melt-infiltrated SiC–Mo5Si3C1 composites. J Am Ceram Soc 1999;82(8):2276–8. [6] Zhu Q, Shobu K. High-temperature fracture toughness of SiC–Mo5Si3C composites. J Mater Sci Lett 2000;19:1529–31. [7] Zhu Q, Shobu K. High-temperature mechanical properties of SiC–Mo5(Si,Al)3C composites. J Am Ceram Soc 2001;84(2):413–9. [8] Nowotny H, Parthe E, Kieffer R, Benesovsky F. The three-component system Mo–Si–C. Monatsh Chem 1954;85:255–72. [9] Buchheit A, Fahrenholtz W, Hilmas G, Deason D, Wang H. Thermal properties of a hot pressed Mo5Si3C1 ceramic, Intermetallics, in press. [10] North B, Gilchrist K. Effect of impurity doping on a reaction-bonded silicon carbide. Ceram Bull 1981;60(5):549–54. [11] Keppler U. Thermal expansion of silicon. Z Metallkd 1988;79(3):157–8. [12] van Loo F, Smet F, Rieck G, Gerard D, Verspui G. Phase relations and diffusion paths in the molybdenum–silicon–carbon system at 1200 C. High Temp. – High Pressures 1982;14(1):25–31. [13] Martinelli A, Drew R, Berriche R. Correlation between the strength of SiC–Mo diffusion couples and the mechanical properties of the interfacial reaction products. J Mater Sci Lett 1996;15:307–10. [14] Morozumi S, Kikuchi M, Sugai S, Hayahsi M. Reaction between molybdenum and carbon, and several carbides. Nippon Kinzoku Gakkaishi 1980;44(12): 1404–13. [15] Suzuki Y, Niihara K. Synthesis and mechanical properties of Mo5Si3C1 and Mo5Si3C1-based composites. Intermetallics 1998;6:7–13. [16] Turan S, Knowles KM. a / b Reverse phase transformation in silicon carbide in silicon nitride–particulate-reinforced-silicon carbide composites. J Am Ceram Soc 1996;79(11):2892–6. [17] Kieffer AR, Ettmayer P, Gugel E, Schmidt A. In: Henisch HR, Roy R, editors. Silicon carbide-1968. Elmsford, NY: Pergamon; 1969. p. S153–66. [18] Turan S, Knowles KM. A comparison of the microstructure of silicon nitride– silicon carbide composites made with and without deoxidized starting material. J Microsc 1995;177(3):287–304. [19] Kasap SO. Principles of electronic materials and devices. 3rd ed. Boston: McGraw Hill; 2006. [20] Costa e Silva A, Kaufman M. Phase relations in the Mo–Si–C system relevant to the processing of MoSi2–SiC composites. Metall Mater Trans A 1994;25A(1): 5–15. [21] Kingery WD, Bowen HK, Uhlmann DR. Introduction to ceramics. New York: John Wiley & Sons; 1976.