Effects of initial α-phase content on the electrical, thermal, and mechanical properties of pressureless solid-state sintered SiC ceramics


 αand β-SiC starting powders of similar particle sizes were used to investigate the effect of initial α-phase content on the electrical, thermal, and mechanical properties of pressureless solid-state sintered (PSS) SiC ceramics with B4C and C. For β-SiC starting powders, a coarse-grained microstructure with elongated platelet grains formed by the 3C to 6H to 4H-SiC phase transformation was obtained. In contrast, α-SiC powders exhibited a fine-grained microstructure with platelet grains. The electrical resistivity decreased by an order of magnitude with increasing initial α-phase content presumably due to (1) an increased 6H-SiC content causing a decrease in bandgap energy and (2) the low soluble impurity content (Fe and V) of the α-SiC powders. The thermal conductivity increased by approximately 32% with increasing initial α-phase content due to (1) an increased 6H-SiC content, which has a higher intrinsic thermal conductivity compared to 4H and (2) the low impurity content of the α-SiC powders. The flexural strength increased by approximately 16% with increasing initial α-phase content due to a decreased flaw size with decreasing grain size. However, the fracture toughness and hardness were insensitive to the change in initial α-phase content.


Introduction
Pressureless solid-state sintered (PSS) silicon carbide (SiC) ceramics exhibit superior hardness, oxidation resistance, corrosion resistance, and wear resistance compared to those of liquid phase sintered (LPS) SiC ceramics . However, there has been limited research on PSS SiC ceramics presumably due to their moderate mechanical properties and high processing temperatures compared to those of LPS SiC ceramics. Recently, there is renewed interest in PSS SiC ceramics due to their desirable engineering properties, which are well suited for various industrial applications such as wafer boats, susceptors, focus rings, and dummy wafers in the semiconductor processing industry; and mechanical seals and bearings in uid machinery.
Prochazka [2] rst developed PSS SiC ceramics using B and C as sintering aids. It was postulated that B segregates at the grain boundaries and reduces the grain boundary energy (γ GB ) and C increases the surface energy (γ SV ) of SiC by removing the native SiO 2 lm via a carbothermal reduction process. The combined effect of both additives reduces γ GB /γ SV to a subcritical value, thus making the densi cation process thermodynamically feasible [23]. Unlike LPS ceramics [24][25][26][27][28][29][30][31][32][33][34][35][36][37][38][39], where mechanical, thermal, and electrical properties can be tuned using a wide range of additive compositions, the scope of PSS SiC ceramics is constrained by the limited choice of sintering additives. The additive systems are limited to elemental B and B-compounds such as B 4 C powders, BCl 3 gas, and LiBH 4 and H 3 BO 3 solutions [1].
To harness the full potential of PSS SiC ceramics, it is important to explore new additives and processing strategies. Recently, BN has been reported as a new sintering additive for PSS SiC ceramics [40]. One potential strategy is to use different crystalline phases (α/β) of the starting SiC powders. Previous works [24,25] on LPS SiC ceramics had established that microstructure and mechanical properties could be tailored by varying the initial α-and β-phase content. It was demonstrated that increasing initial β-phase content was bene cial in improving the fracture toughness attributed to the crack de ection by α-SiC platelet grains formed in situ via β to α phase transformation during sintering [24,25] whereas a negrained microstructure with equiaxed grains resulted in improved exural strength with increasing initial α-phase content [24]. For PSS SiC ceramics, a previous work [7] revealed interesting microstructural characteristics for α-SiC and β-SiC powders sintered with 0.4 wt% B 4 C and 2 wt% C at 2150°C in owing Ar gas. For α-SiC powders, a ne-grained microstructure was obtained. On the other hand, β-SiC starting powders exhibited a coarse-grained microstructure with elongated platelet grains of high aspect ratio formed via β to α phase transformation of SiC during sintering. However, there was little systematic research to be found on the effects of initial α-SiC content on the electrical, thermal, and mechanical properties of PSS SiC ceramics.
This study is a novel investigation on the effect of initial α-SiC content (0-100%) on the electrical, thermal, and mechanical properties of PSS SiC ceramics. Given that the initial α-and β-phase content have a profound effect on the microstructure [7], it would be interesting to investigate the effect of such tailored microstructures on the electrical, thermal, and mechanical properties of PSS SiC ceramics.

Experimental
Commercially available α-SiC (0.5 µm, Sintex 15C, Sika Tech, Fiven, Norway), β-SiC (~ 0.5 µm, BF17, H.C. Starck, Germany), B 4 C (0.5 µm, Kojundo Chemical Co., Ltd., Japan), and phenolic resin (TD-739, Kangnam Chemical Co., Ltd., Korea) were used as the starting materials. As shown in Table 1, the impurity contents obtained using inductively coupled plasma-optical emission spectroscopy (ICP-OES, OPTIMA 8300 DV, Perkin Elmer, USA) revealed a higher concentration of Fe and V impurities in β-SiC powders compared to α-SiC powders. The phenolic resin had a carbon yield of 60%. As shown in Table 2, six batches of powders with varying initial α-SiC content (0-100% of total SiC) were prepared. The specimens were designated as Ax, where 'x' denotes the wt% of α-SiC phase in the total SiC content. B 4 C and C contents were xed at 0.7 and 2.5 wt%, respectively, in all batches. The batches were mixed using SiC balls by ball milling in ethanol with a ball to powder weight ratio of 2:1. After drying, the powder mixture was granulated using a 60-mesh sieve. The sieved powders were uniaxially pressed at 25 MPa into both cylindrical (Ø 25 mm × 7 mm high) and rectangular hexahedron (40 × 40 × 4 mm 3 ) compacts, followed by cold isostatic pressing at 207 MPa. The green compacts were pyrolyzed at 900°C in owing Ar gas for 1 h and subsequently pressureless sintered in a graphite resistance furnace at 2100°C for 2 h in owing Ar gas. The theoretical and sintered densities were determined by the rule of mixtures and the weight-tovolume ratio of the sintered specimens, respectively. The polished specimens (up to 1 µm) were etched using a modi ed Murakami's reagent [41]. The details of microstructure characterization using SEM, phase analysis using XRD, electrical properties measurement using Hall effect, thermal conductivity measurement using laser ash method, exural strength measurement using ASTM C 1161-13, fracture toughness measurement using JIS-R1607, and hardness measurement using Vickers indentation method are described elsewhere [35,39,42,43]. Impurity contents of α-and β-SiC starting powders obtained using inductively coupled plasma -optical emission spectroscopy (ICP-OES).

Phase analysis and microstructure
As shown in Table 2, all specimens had high relative densities (≥ 99%) when pressureless solid-state sintered at 2100°C for 2 h in owing Ar gas using 0.7 wt% B 4 C and 2.5 wt% C as sintering aids. The XRD patterns and SiC polytype contents in the sintered specimens are shown in Fig. 1 and Table 2, respectively. The XRD patterns revealed α-SiC (6H, 4H, 15R polytypes), and traces of B 4 C and graphite in the sintered specimens. β-SiC (3C) was not detected in the sintered specimens because of its complete conversion to α-phase caused by B-doping in the SiC lattice [38] and the high sintering temperature [44]. Residual C left after the carbothermal reduction of native SiO 2 was transformed to graphite at the sintering temperature [5]. As shown in Table 2, 4H-SiC content decreased with an increase in initial αphase content due to a slower 6 H to 4H-SiC phase transformation in α-SiC powders compared to β-SiC powders [7,[24][25][26]. The 3C to 6H to 4H phase transformation was accelerated in the β-SiC powder due to its high thermodynamic instability at the sintering temperature and high Al content (refer to Table 2) [45][46][47].
Typical microstructures of PSS SiC ceramics with varying initial α-phase content are shown in Fig. 2. The SEM micrographs exhibited a coarse-grained microstructure with elongated α-SiC platelet grains formed in situ via an accelerated 3C to 6H to 4H phase transformation in A0 and a gradual transformation to ne-grained microstructure with increasing initial α-phase content. The average aspect ratio of SiC grains decreased from 7.4 ± 1.3 in A0 to 2.9 ± 0.4 in A100. The ne-grained microstructure was responsible for the increased density (99.1→99.8%) with increasing initial α-phase content (refer to Table 2). It is well documented that equiaxed grains can easily accommodate each other and annihilate the residual pores, a phenomenon known as grain-shape-accommodation [29]. On the contrary, platelet grains grown preferentially along the c-axis hinder the elimination of inter-grain pores, due to interlocked elongated grains, resulting in a lower relative density.

Electrical and thermal properties
The electrical resistivity and carrier density of PSS SiC ceramics with varying initial α-phase content are shown in Fig. 3. The electrical resistivity monotonically decreased (2.6 × 10 7 → 1.5 × 10 6 Ω cm) with increasing initial α-phase content presumably due to (1) a decreased bandgap energy with increasing 6H-, 15R-and decreasing 4H-SiC contents (refer to Table 2) and (2) the low impurity content of the α-SiC powders ( Table 1). The high content of Fe and V impurities in β-SiC powders resulted in the donoracceptor compensation and trapping of charge carriers (refer to Table 1) [48], which is also indicated by an increase in carrier density (1.1 × 10 10 → 2.8 × 10 13 cm − 3 ) with increasing initial α-phase content (refer to Fig. 3). The bandgap energies of 6H, 15R, and 4H, are 3.10, 2.98, and 3.29 eV, respectively [49,50]. Various metallic and non-metallic impurities are known to affect the electrical properties of SiC ceramics [48,51]. Impurities like Al, B, Ga, and Sc are p-type dopants, whereas N, P, and Fe are n-type dopants. Both n-and p-type dopants increase the charge carrier density and thus decrease the electrical resistivity of SiC ceramics. However, in the case of codoping of n-and p-type dopants, an acceptor-donor compensation takes place causing a decrease in the charge carrier density [52,53]. On the other hand, impurities like Be, O, V tend to create a deep level within the bandgap trapping the charge carriers causing an increased electrical resistivity [53]. Table 3, the electrical resistivity of the solid-state sintered SiC ceramics could be tuned in the range of 10 1 -10 11 Ω cm by changing the composition and sintering atmosphere [14,17,21,54]. For example, PSS SiC ceramics with high electrical resistivities of 10 8 -10 11 Ω•cm were fabricated by sintering in an N 2 atmosphere at ≥ 2275°C and the high electrical resistivity was attributed to the formation of an electrically insulating C-B-N intergranular phase at grain boundaries [54]. On the other hand, the electrical resistivity of PSS SiC ceramics with 1 wt% B 4 C could be decreased (10 6 →10 2 Ω cm) by four orders of magnitude by adding 2 wt% electrically conductive graphene [21]. The ceramics in this study exhibited intermediate electrical resistivities (10 6 -10 7 Ω•cm) comparable to those of PSS SiC ceramics sintered with 1 wt% B 4 C at 2050°C in Ar gas [21]. As shown in Table 3, the reports on the electrical resistivity of PSS SiC ceramic sintered with β-SiC powders are rarely found, presumably due to their higher cost and lower stability at high sintering temperatures compared to α-SiC powders. This study revealed that the high thermodynamic instability of β-SiC powders can be used to tune the electrical properties of PSS SiC ceramics to some extent via a controlled 3C to 6H to 4H phase transformation. Staring powder polytype, processing conditions, relative density, electrical resistivity, and thermal conductivity of solid-state sintered SiC ceramics.

As shown in
The thermal conductivity of PSS SiC ceramics with varying initial α-phase content is shown in Fig. 4. The thermal conductivity monotonically increased from 117.1 Wm − 1 K − 1 in A0 to 154.2 Wm − 1 K − 1 in A100 due to (1) an increased 6H-SiC content, which has a higher intrinsic thermal conductivity compared to that of 4H-SiC [38] and (2) the low soluble impurity content of the α-SiC starting powders (refer to Table 1). Soluble impurities, such as Fe and V [55,56], act as phonon scattering sites in the SiC lattice and deteriorate the thermal conductivity. As shown in Table 3, the reported thermal conductivity of solid-state sintered SiC ceramics varies from 107 to 192 Wm − 1 K − 1 depending on the composition and processing conditions. Li et al. [12] reported a high thermal conductivity of 192.2 Wm − 1 K − 1 attributed to an optimum content of sintering aids (0.4 wt% B 4 C and 3 wt% C) derived from polymeric precursors. However, a low thermal conductivity of 107 Wm − 1 K − 1 was reported for SiC-5 wt% graphene composite attributed to an increased interfacial thermal resistance at SiC-graphene heterogeneous interfaces [14]. As shown in Table 3, the literature is limited to the thermal conductivity of PSS SiC ceramics started with α-SiC powders. This study revealed that β-SiC powders are detrimental to thermal conductivity because they promote 6H to 4H phase transformation, leading to an increase of 4H phase, which has a low intrinsic thermal conductivity compared to 6H phase.
Based on the above discussion, the thermal conductivity of PSS SiC ceramics can be improved by (1) minimizing the content of sintering additives and secondary phases, (2) using high purity starting powders, and (3) avoiding or minimizing 6H to 4H phase transformation during sintering.

Mechanical properties
The mechanical properties of PSS SiC ceramics with varying initial α-phase content are shown in Table 4. The exural strength (σ f ) increased from 289.5 MPa in A0 to 335.1 MPa in A100 attributed to a decreased defect size with decreasing grain size [11,13,24,57]. However, the fracture toughness (K IC ) and hardness were almost constant at 2.2-2.3 MPa m 1/2 and approximately 28 GPa, respectively, with the variation in initial α-SiC content. As shown in Fig. 5, the crack path did not exhibit any active crack-microstructure interaction due to strong interfacial bonding between the SiC grains, leading to a predominant transgranular fracture in all specimens (refer to Fig. 6). In the absence of active crack-microstructure interaction, the fracture toughness was insensitive to changes in microstructure. These results are consistent with the ndings of Magnani et al. [11] (refer to Table 5). It was demonstrated that ne-grained PSS SiC ceramics fabricated via two-step sintering exhibited 39% higher exural strength compared to the coarse-grained ceramics fabricated via a conventional sintering process [11]. However, the fracture toughness and hardness did not exhibit any signi cant change [11]. As shown in Table 4, the calculated critical aw sizes (c, 26-35 µm) were larger than the average grain sizes implying that the fracture had originated from processing defects. The critical aw sizes were calculated using the equation K IC = 1.35 σ f c 1/2 , assuming semielliptical surface aws [24]. Figure 7 shows the typical fracture origins of PSS SiC ceramics which were mostly residual pores near the surface. The critical aw sizes of A0 and A100 observed in Fig. 7 are comparable to the sizes calculated (Table 4). A decreased critical aw size with increasing initial α-phase content indicates a superior sinterability of α-SiC compared to β-SiC powders, which is also indicated by a decreased scatter in exural strength values with the addition of α-SiC powders (refer to Table 4).   [11,13] reported high exural strengths of 530-556 MPa for two-step sintered SiC ceramics with ne-grained microstructure compared to those (320-467 MPa) of PSS SiC ceramics with coarse-grained microstructures reported in other works [3,4,6,15,16]. A ne-grained microstructure is bene cial in improving exural strength because of the reduced critical aw size [11,13]. As shown in Table 5

Conclusions
SiC ceramics with 0 to 100% initial α-phase content were densi ed to high relative densities (≥ 99%) via pressureless solid-state sintering with 0.7 wt% B 4 C and 2.5 wt% C at 2100°C in owing Ar gas. The major ndings are summarized as follows: 1. The phase analysis revealed an increased 6H-SiC content with increasing initial α-phase due to a slower 6H to 4H-SiC phase transformation in α-SiC powders compared to β-SiC powders. 2. α-SiC starting powders exhibited a ne-grained microstructure with platelet grains of low aspect ratio (2.9) whereas β-SiC powders exhibited a coarse-grained microstructure with elongated platelet grains of high aspect ratio (7.4), which formed in situ via an accelerated 3C to 6H to 4H phase transformation of SiC.
3. The electrical resistivity monotonically decreased from 2.6 × 10 7 to 1.5 × 10 6 Ω cm with increasing initial α-phase content from 0 to 100% attributed to (1) an increased 6H-SiC content causing a decrease in bandgap energy and (2) the low soluble impurity content (Fe and V) of the α-SiC powders. The high content of Fe and V impurities in β-SiC powders resulted in the donor-acceptor compensation and trapping of charge carriers.
4. The thermal conductivity monotonically increased from 117.1 to 154.2 Wm -1 K -1 with increasing initial α-phase content from 0 to 100% due to (1) an increased 6H-SiC content with a high intrinsic thermal conductivity and (2) the low impurity content of the α-SiC powders.

Figure 3
Electrical resistivity and carrier density of pressureless solid-state sintered SiC ceramics with varying initial α-phase content.

Figure 4
Thermal conductivity of pressureless solid-state sintered SiC ceramics with varying initial α-phase content.