Low-Temperature Pressureless Sintering of a-SiC with Al 4C3–B4C–C Additions

Journal
J. Am. Ceram. Soc., 82 [8] 1959–64 (1999)
Low-Temperature Pressureless Sintering of ?-SiC with Al4C3–B4C–C Additions
You Zhou,*,? Hidehiko Tanaka,* Shigeki Otani, and Yoshio Bando
National Institute for Research in Inorganic Materials, Tsukuba, Ibaraki, 305-0044 Japan
Various combinations of Al4C3, B4C, and carbon were used as sintering aids for pressureless sintering of ?-SiC (94% 6H + 6% 15R) powders. Densification behavior, polytypic transformation, microstructural development, and mechanical properties were studied. Appropriate amounts of Al4C3–B4C–C additions could facilitate ?-SiC powders being pressureless sintered to high densities (>95% TD) at temperatures ≥1850°C. Even using an ?-SiC as a starting powder, many elongated grains developed in the microstructures during sintering at low temperature that might contribute to toughening. Anisotropic grain growth was found to be associated with the faulted structures and the 6H → 4H phase transformation. I. Introduction
hot-pressing, which exhibited high fracture toughness and strength. Inomata et al.19 investigated the phase relation in a SiC– Al4C3–B4C system and found that a liquid phase was formed around Al8B4C7 at 1800°C. In the present work, various combinations of Al4C3, B4C, and carbon were used as sintering aids for the pressureless sintering of ?-SiC with the aim of lowering the sintering temperature. The process of successful lowtemperature pressureless sintering of ?-SiC may provide a means of achieving the efficient, economical, and large-scale production of dense SiC sintered bodies with complex shapes and good mechanical properties. II. Experimental Procedure The starting powder was ?-SiC (Showadenko Co., Tokyo, Japan) with an average particle size of 0.48 ?m. Purity of the SiC was >99%, and the contents of oxygen and free carbon were 0.18 and 0.55 wt%, respectively. The polytype composition of the SiC powder was 94% 6H plus 6% 15R, which was determined by XRD method.20 High-purity Al4C3 powder (High Purity Chemicals Co., Saitama, Japan) and B4C powder (Cerac Inc., Milwaukee, WI) were used as additives. The third additive, carbon, was introduced by carbonization of phenolic resin (Dainipponn Ink & Chemicals, Inc., Tokyo, Japan), which resulted in 1.8 wt% carbon of the total weight of SiC + Al4C3 + B4C. The ?-SiC, Al4C3, B4C, and phenolic resin were wet blended in ethanol into various compositions (see Table I) along with 2 wt% poly(ethylene glycol) as a binder. Mixing was performed in a planetary ball mill, with all balls and the pot consisting of SiC ceramic material. After mixing for 8 h, the slurry was dried at 110°C in an oven for 5 h, then the mixture was crushed in a mortar with a pestle and screened through a 100 mesh sieve. Green bars, ～40.0 mm × 5.4 mm × 4.8 mm, were formed by uniaxial pressing at a pressure of 20 MPa, followed by isostatic pressing at a pressure of 200 MPa. Sintering was performed in a carbon resistance furnace. The furnace was heated at a rate of 50°C/min from room temperature to 1500°C under vacuum and held for 30 min to remove SiO2 from the surface of SiC.21 Then, pure argon was introduced, and the temperature was raised at a rate of 8°C/min to the desired sintering maxima (1800–2000°C) and held for 1 h or 30 min. Holding time was 30 min only when the sintering temperature was 2000°C. The densities of the ceramics were determined by the Archimedes method. Sintered bodies were ground into powders with a WC mill, and XRD with CuK? radiation was used to
Table I.
Sample No.
covalency of Si–C bonds and the low self-diffusion coefficient make it impossible to sinter SiC to high densities without additives or high external pressure.1 A great deal of research has been conducted to find appropriate sintering additives for SiC powders. In the early 1970s, Prochazka2 first achieved pressureless solid-state sintering of SiC by using additions of boron and carbon to submicrometer ?-SiC powders. Later, ?-SiC powder was also densified with the aid of boron and carbon.3 Although the pressureless sintering of SiC containing boroncarbon additives has become a key technology in manufacturing SiC sintered bodies for applications as mechanical components, there are still some problems with this process, e.g., low fracture toughness of the materials and processing temperatures >2100°C that easily result in exaggerated grain growth.4 Another approach to the densification of SiC is liquid-phase sintering. Omori and Takei5–7 found that a wide variety of rare-earth oxides, usually in combination with Al2O3 and/or boron compounds, could be used to promote the densification of SiC via a liquid phase formed during the heating process. So far, Al2O3,8–10 Al2O3–Y2O3 (YAG),11–13 and Al2O3–Y2O3– CaO14 have also been shown to be effective sintering aids for SiC. Besides the above-mentioned oxides, aluminum-boroncarbon nonoxide system additives can also serve as sintering aids for liquid-phase sintering of SiC. Shinozaki et al.15,16 studied the pressureless sintering of ?-SiC with aluminum, boron, and carbon additions and found that the addition of aluminum not only initiated densification at substantially lower temperatures than boron and carbon alone but also enhanced the ?-to?-phase transformation. Recently, Cao et al.17,18 also used aluminum, boron, and carbon as the additives and ?-SiC as the starting powder to produce in situ toughened SiC ceramics by
HE HIGH
T
Compositions and Theoretical Densities of Various Mixtures
Composition (wt%) ?-SiC Al4C3 B4C C TD (g/cm3)
Molar ratio of Al4C3/B4C
D. K. Kim—contributing editor
Manuscript No. 190228. Received April 28, 1998; approved December 3, 1998. *Member, American Ceramic Society. ? Present address: National Industrial Research Institute of Nagoya, Nagoya, Japan.
A01B4 A13B4 A12B4 A11B4 A21B4 A31B4 1959
0 1/3 1/2 1/1 2/1 3/1
97.83 97.50 97.33 96.82 95.80 94.78
0.00 0.34 0.51 1.02 2.04 3.06
0.39 0.39 0.39 0.39 0.39 0.39
1.77 1.77 1.77 1.77 1.77 1.77
3.19 3.19 3.19 3.18 3.17 3.16

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analyze the crystallographic composition and polytype content of the ceramics according to Tanaka’s method.20 The polishing surfaces of the sintered bars were etched by Murakami’s reagent (10 g of NaOH and 10 g of K3Fe(CN)6 in 100 mL H2O at 110°C), and SEM coupled with energy dispersive X-ray spectroscopy (EDX) were used to observe the morphology and analyze the chemical composition. Microstructural characterization of the ceramics was further performed by TEM, selected area diffraction (SAD), and high-resolution transmission electron microscopy (HREM). A foil was prepared by standard techniques involving cutting, grinding, dimpling, argon-ion thinning to perforation, and coating with a thin carbon film. Fracture toughness was measured by the indentation fracture (IF) method according to Japanese Industrial Standard JIS R-1607.22 Ten indentations were made at 49 N load on the polished surface of each sintered sample. The elastic modulus of every ceramic sample, which was needed to calculate fracture toughness in the IF method, was measured by the pulse echo method. III. Results and Discussion
(1) Densification Behavior According to Inomata et al.,19 who reported that a liquid phase formed around Al8B4C7 (in which the molar ratio of Al4C3/B4C was 2) at 1800°C in the SiC–Al4C3–B4C system, various amounts of Al4C3–B4C–C additives were used for the purpose of investigating the effect of liquid phase on the sintering behavior of ?-SiC powders. For comparison, a composition with B4C–C additive (designated as A01B4 in Table I) was also prepared. The total amount of additives (Al4C3 + B4C + C) was varied by changing the molar ratio of Al4C3/B4C while keeping the amount of B4C addition as a constant (0.39 wt%) (designated as A13B4, A12B4, A11B4, A21B4, and A31B4 in Table I). The amount of carbon addition was set to be 1.77 wt%. Figure 1 shows the densities of the sintered SiC samples containing various amounts of additives as a function of sintering temperature. Combinations of Al4C3, B4C, and carbon as sintering additives were very effective in promoting the sintering of ?-SiC powders. For samples A11B4 and A21B4, the
densities were already near 91% TD, while the sintering temperature was ?1800°C. At 1850°C, the samples were sintered to high densities (>95% TD). This sintering temperature was ～250°C lower than that required for normal solid-state sintering of SiC powders with boron-carbon additives to attain such high densities. The sintering of the SiC powders with Al4C3– B4C–C additions was thought to be a liquid-phase sintering process (as confirmed by the following experimental results), similar to that for aluminum-, boron-, and carbon-doped SiC described by Cao et al.17 As measured by powder XRD analysis, diffraction peaks having lattice spacings of d ? ??0.291, 0.237, 0.171 nm, etc., were detected in samples with Al4C3– B4C–C additions sintered at 1800°C. These diffraction peaks corresponded to those with Al8B4C7, whose melting point was near 1800°C. During sintering, Al4C3 and B4C seemed to react and form an aluminum boron carbide compound (Al8B4C7), which was a liquid phase at ～1800°C. Figure 1 also shows that, for samples doped with different amounts of additives, the effect of sintering temperature on densification was different. For samples containing relatively small amounts of additives (e.g., A12B4 and A11B4, whose total quantity of Al4C3 and B4C was <2 wt%), the densities of the sintered bodies increased with increased sintering temperature (between 1850° and 2000°C). However, for samples containing relatively large amounts of additives (e.g., A21B4 and A31B4, whose total quantity of Al4C3 and B4C was >2 wt%), the densities of the ceramics hardly changed as sintering temperature increased from 1850° to 2000°C due to the effect of grain growth. In the present case, the liquid phase formed in the powder compacts during heating was, at most, only several percent by volume, so the contribution of the particle rearrangement mechanism to densification was very small.23 The dominant sintering mechanism might have been solution– reprecipitation, and fast mass transfer through the liquid resulted in quick densification at lower temperatures compared with solid-state sintering. However, during such a liquid-phase sintering process, especially in the intermediate and final stages, densification always occurred in conjunction with microstructural coarsening, because solution–reprecipitation contributed simultaneously to both grain growth and densification. This was confirmed by microstructural observations (Figs. 2 and 3).
Fig. 1.
Variation of density with sintering temperature for SiC samples doped with various amounts of additives.

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Inomata et al.19 indicated that Al8B4C7 was the only composition of the liquid phase formed through the chemical reaction 2Al4C3 + B4C → Al8B4C7 while heating mixtures of Al4C3 and B4C; this was also confirmed by our XRD analysis. Then, the compositions of the liquid phases formed in samples A13B4, A12B4, A11B4, A21B4, and A31B4 might be considered approximately the same, i.e., Al8B4C7, but the quantities of liquid phases in the different samples were different. In the present sintering system, a small amount of liquid was indispensable in achieving low-temperature densification (e.g., A13B4, A12B4, and A11B4). When the amount of liquid was too high (e.g., A21B4 and A31B4), grain growth would prevail over densification; i.e., before the compact reached a high density, the grains had already grown rapidly and formed a rigid skeleton that retarded further densification. In such a case, increasing sintering temperatures were not effective in improving densification but rather accelerated grain growth. So, optimum values of the Al4C3–B4C–C additives existed for achieving successful low-temperature sintering of ?-SiC powders. When the content of the additives was too high, the sample could not be sintered to high density at any temperature. When the content of the additives was too low, the sample could not be sintered to high density at low temperatures (e.g., 1850°C), but it could be completely densified at higher temperatures (e.g., 2000°C), which were still rather low compared with those of the solid-state sintering process. (2) SEM Study and Crystallographic Analysis Figures 2 and 3 show the SEM images of the polished and etched surfaces of the samples doped with various amounts of additives sintered at 1850° and 2000°C, respectively. As shown in Fig. 2, after sintering at 1850°C, sample A01B4 (doped with 0.39 wt% B4C and 1.77 wt% C) was still porous and did not undergo obvious grain growth. For sample A13B4 with addition of 0.34 wt% Al4C3 together with B4C and carbon, the microstructure was denser, grain sizes were bigger, and the number of fine particles were reduced. With greater additions of Al4C3, sample A12B4 underwent more obvious grain growth, and samples A11B4, A21B4, and A31B4 exhibited very coarse microstructures. A very interesting phenomenon was that there were many elongated grains in the microstruc-
tures of samples A11B4, A21B4, and A31B4; i.e., a monolithic SiC ceramic body containing in-situ-formed elongated grains could be fabricated at low sintering temperatures by using appropriate sintering additives with ?-SiC powder as the starting material. Even though exaggerated grain growth of SiC at >2100°C was usually anisotropic because of the large differences in the interfacial energies of different crystallographic planes, the microstructures of SiC ceramics made by sintering ?-SiC starting powders at ?1850°C were always composed of equiaxial grains. ?-SiC powder was always used as the starting material to fabricate SiC ceramics with microstructures consisting of elongated grains via either the ? → ? phase transformation12 or seeding effect24 during a subsequent annealing process. As shown in Fig. 3, after sintering at 2000°C, the microstructure of sample A01B4 was still fine, and the grains were all equiaxial. Whereas all other samples doped with Al4C3– B4C–C in their starting composition underwent enormous anisotropic grain growth, both the grain sizes and relative content of the elongated grains increased with an increase in the amount of additives (Figs. 3(b)–(f)). For samples A21B4 and A31B4, some large pores still existed in the microstructures (Figs. 3(e) and (f)), and their densities did not surpass those of the corresponding A21B4 and A31B4 samples sintered at 1850°C, as shown in Fig. 1. The reason could be that a relatively large amount of liquid formed in sample A21B4 or A31B4 during heating and facilitated very fast grain growth, and the large elongated grains formed a skeletal microstructure that provided rigidity to the compact and, thereby, inhibited further pore elimination. The crystallographic transformation behavior of the various samples during sintering was studied, and the polytype compositions are listed in Table II. As shown, when sintered at temperatures between 1850° to 2000°C, no polytypic transformation occurred in sample A01B4. In contrast, the 6H → 4H polytypic transformation took place in all samples with Al4C3– B4C–C additions, and the transformation proceeded to a higher degree with an increase in the amount of additives. This polytypic transformation might have contributed to anisotropic grain growth, which is discussed further in the following section.
Fig. 2. SEM of the chemically etched surfaces of samples sintered at 1850°C: (a) A01B4, (b) A13B4, (c) A12B4, (d) A11B4, (e) A21B4, and (f) A31B4.

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Fig. 3. SEM of the chemically etched surfaces of samples sintered at 2000°C: (a) A01B4, (b) A13B4, (c) A12B4, (d) A11B4, (e) A21B4, and (f) A31B4.
Table II.
Sample No.
Contents of Polytypes in Various SiC Ceramics?
Sintering temperature (°C) Polytype content (%) 4H 6H 15R
A01B4
A13B4
A12B4
A11B4
A21B4
1850 1900 1950 2000 1850 1900 1950 2000 1850 1900 1950 2000 1850 1900 1950 2000 1850 1900 1950 2000
0 0 0 0 1 2 3 4 2 2 3 4 2 6 7 10 6 10 14 18
94 94 94 94 93 92 92 91 92 92 92 91 92 89 89 86 89 86 82 77
6 6 6 6 6 6 5 5 6 6 5 5 6 5 4 4 5 4 4 5
?
Polytype composition of the starting ?-SiC powder is 94% 6H + 6% 15R.
EDX was used to conduct a chemical analysis to detect the distribution of the sintering aids (mainly aluminum) in the microstructure of ceramics made by sintering ?-SiC with Al4C3–B4C–C additions. Most of the aluminum resided at the edge (surface or grain boundary) of the SiC grains, while a very small amount of aluminum entered into the SiC lattice. (3) TEM Study TEM analyses were conducted on ceramics obtained by sintering ?-SiC powders with Al4C3–B4C–C additions. TEM observations showed that most grains exhibited heavily faulted structures. Furthermore, SAD analyses showed that, in the microstructures, most grains were faulted 6H- or 4H-SiC structures and a few were perfect 6H or 4H structures. A partially faulted 6H-SiC grain is shown in Fig. 4. In the SAD patterns of many of the grains, continuous streaks frequently occurred along the h0.1 reciprocal lattice rows (Fig. 4(c)). This streaking was caused by one-dimensionally disordered stacking of the close-packed SiC lattice planes.25 Some reports have claimed that an aluminum impurity favored the formation of the 4H polytype during the ? → ? SiC transformation.26,27 In the present case, the aluminum atoms in the SiC lattice might have initiated and enhanced the formation of the stacking faults.
Fig. 4. (a) TEM image of a partially faulted SiC grain in sample A11B4 sintered at 1900°C; (b) [1120] diffraction pattern from area “A” (in (a)) showing a perfect 6H-SiC structure; and (c) [1120] diffraction pattern from area “B” (in (a)) showing continuous fine streaks connecting the 6H reflections along the h0.1 reciprocal lattice rows, indicative of a faulted 6H-SiC structure.

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Fig. 5. (a) TEM image of a partially transformed SiC grain in sample A11B4 sintered at 1900°C; (b) lattice fringe image of one part of the grain; (c) [1120] diffraction pattern from the right part in (b) showing a perfect 4H-SiC structure; (d) SAD pattern from the left part in (b) showing the overlapping [1120] diffraction spots of 6H- and 4H-SiC, together with continuous fine streaks resulting from stacking faults; and (e) indexed diagram of the diffraction patterns in (d).
Table III.
Sintering temperature (°C) A01B4
Fracture Toughness of Various Ceramics
Fracture toughness of the sintered samples (MPaиm1/2)? A13B4 A12B4 A11B4 A21B4
1850 1900 1950 2000
?
2.7 ± 0.1
3.5 ± 0.1 3.0 ± 0.1 2.8 ± 0.1
3.5 ± 0.1 3.1 ± 0.2 2.8 ± 0.2
3.9 ± 0.2 3.6 ± 0.2 3.5 ± 0.1 3.4 ± 0.1
4.3 ± 0.1 4.2 ± 0.3 3.5 ± 0.1 3.5 ± 0.1
Fracture toughness of ceramics with relative densities <95% TD is not measured, because IF method is not suitable for low-density ceramic samples.
These stacking faults might have served as nucleation sites for the 6H → 4H transformation, similar to the ? → ? SiC transformation starting from heavily faulted ?-polytype, reported by Heuer et al.28,29 Figure 5 shows a partially transformed SiC grain composed of 4H and 6H polytypes, together with a few stacking faults. It is thought that the 4H polytype resulted from the 6H → 4H transformation. The lattice fringe image also revealed a (0001)4H?(0001)6H relationship between the newly formed 4H and the untransformed 6H, which led to an overlapping of the [1120] diffraction spots of 6H and 4H. The 6H → 4H polytype transformation might have played an important role in initiating anisotropic grain growth, which resulted in the formation of elongated SiC grains in the sintered bodies, similar to the 3C → 4H transformation that was associated with the grain growth of platelike grains during the hot-pressing of ?-SiC as reported by Cao et al.17 (4) Fracture Toughness Table III lists the fracture toughness of the SiC ceramic samples doped with various amounts of additives and sintered at various temperatures. For samples sintered at the same tem-
peratures, the fracture toughness values showed a tendency to increase with an increase in the amount of additives. That could be because more additives resulted in more elongated SiC grains with higher aspect ratio (Figs. 2 and 3), and the large elongated grains were usually effective in increasing the fracture toughness of the ceramics by crack deflection and/or crack bridging, as proved by many researchers.12,18,21,24 So, the Al4C3–B4C–C sintering aids may provide a new way of making self-toughened SiC ceramics at a relatively low sintering temperature from ?-SiC starting powders. IV. Conclusions
Combinations of Al4C3, B4C, and carbon were found to be effective sintering aids for the pressureless sintering of ?-SiC powders. By adding appropriate amounts of such additives, e.g., 0.51–1.02 wt% Al4C3 + 0.39 wt% B4C + 1.77 wt% carbon, ?-SiC powders were pressurelessly sintered to high densities (>95% TD) at temperatures ?1850°C. Liquid-phase sintering using Al4C3–B4C–C additives not only substantially decreased the sintering temperature of ?-SiC but also accelerated grain growth, which was often anisotropic.

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In the microstructures of the ceramics, most grains were heavily faulted. The faulted structures favored the 6H → 4H polytype transformation that might have contributed to anisotropic grain growth. Anisotropic grain growth led to the development of many elongated grains in the microstructure, which tended to increase the fracture toughness of the ceramics.
References
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