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Process-tolerant pressureless-sintered silicon carbide ceramics with alumina-yttria-calcia-strontia
Accepted Manuscript Title: Process-tolerant pressureless-sintered silicon carbide ceramics with alumina-yttria-calcia-strontia Authors: Yu-Kwang Seo, Jung-Hye Eom, Young-Wook Kim PII: DOI: Reference:
S0955-2219(17)30599-X http://dx.doi.org/10.1016/j.jeurceramsoc.2017.09.011 JECS 11446
To appear in:
Journal of the European Ceramic Society
Received date: Revised date: Accepted date:
16-3-2017 9-9-2017 10-9-2017
Please cite this article as: Seo Yu-Kwang, Eom Jung-Hye, Kim YoungWook.Process-tolerant pressureless-sintered silicon carbide ceramics with alumina-yttria-calcia-strontia.Journal of The European Ceramic Society http://dx.doi.org/10.1016/j.jeurceramsoc.2017.09.011 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Process-tolerant pressureless-sintered silicon carbide ceramics with aluminayttria-calcia-strontia Yu-Kwang Seo, Jung-Hye Eom, and Young-Wook Kim* Functional Ceramics Laboratory, Department of Materials Science and Engineering, The University of Seoul, Seoul 02504, Republic of Korea
Abstract Process-tolerant SiC ceramics were prepared by pressureless sintering at 1850-1950oC for 2 h in an argon atmosphere with a new quaternary additive (Al2O3-Y2O3-CaO-SrO). The SiC ceramics can be sintered to a >94% theoretical density at 1800-1950oC by pressureless sintering. Toughened microstructures consisting of relatively large platelet grains and small equiaxed grains were obtained when SiC ceramics were sintered at 1850-1950oC. The presently fabricated SiC ceramics showed little variability of the microstructure and mechanical properties with sintering within the temperature range of 1850-1950oC, demonstrating process-tolerant behavior. The thermal conductivity of the SiC ceramics increased with increasing sintering temperature from 1800oC to 1900oC due to decreases of the lattice oxygen content of the SiC grains and residual porosity. The flexural strength, fracture toughness, and thermal conductivity of the SiC ceramics sintered at 1850-1950oC were in the ranges of 444-457 MPa, 4.9-5.0 MPa∙m1/2, and 76-82 Wm-1K-1, respectively. Keywords: Silicon carbide; Mechanical properties; Thermal conductivity; Microstructure; Process-tolerant behavior *Corresponding author. Fax: +82 2 6490 2404 E-mail address: [email protected] (Y.-W. Kim). 1. Introduction The market for advanced ceramics is continuously growing due to their excellent performance, multifunctional properties, and high energy efficiency [1]. However, the lower yields for ceramics than for those of other materials such as metals and polymers could be a
barrier for widening the industrial applications of advanced ceramics due to cost effectiveness [2,3]. Processing parameters such as the sintering temperature, time, and annealing conditions affect the microstructural characteristics and consequently lead to variability of the electrical, thermal, optical and mechanical properties [4-8]. Variability of these properties can potentially lower the production yield. Thus, finding systems that exhibit “process-tolerant behavior,” i.e., the insensitivity of properties to specific processing parameters, is an important issue in ceramic industries for improving the production yield and cost effectiveness [9]. Silicon carbide (SiC) is one of the most important engineering ceramics due to its good thermal conductivity, good oxidation and corrosion resistance, as well as its excellent mechanical properties [5,7,10-12]. Thus, dense SiC ceramics are widely used in many applications such as heaters, heat exchangers, semiconductor processing parts, mechanical seals, and cladding materials for nuclear fuels. However, the densification of SiC ceramics without sintering additives is difficult due to their highly covalent bonding and low selfdiffusivity. Various sintering additive compositions were investigated for pressurelesssintered SiC ceramics. The relative density and flexural strength of pressureless-sintered SiC ceramics showed wide variations (96–99% and 394–633 MPa, respectively), which depended on the additive composition and processing conditions [13-22]. Previous results include values of 98% and 556 MPa for SiC ceramics sintered with 3.5 wt% B-C [13], 99% and 394 MPa for SiC ceramics sintered with 7 wt% AlN-C [14], 96% and 440 MPa for SiC ceramics sintered with 10 wt% Al2O3-Y2O3-MgO [18], 98% and 433 MPa for SiC ceramics sintered with 3 vol% Al2O3-AlN-Y2O3 [20], 98% and 633 MPa for SiC ceramics sintered with 9 wt% Al2O3-Y2O3-CaO [21], and 97% and 520 MPa for SiC ceramics sintered with 6.5 vol% AlNY2O3-Sc2O3 [22]. However, there is no previous work investigating pressureless-sintered SiC ceramics with a quaternary additive system. The reported pressureless-sintered SiC ceramics
with binary or ternary sintering additives showed a high sintered density and excellent mechanical properties. However, all of the above pressureless-sintered SiC ceramics showed a wide variation of microstructures and mechanical properties depending on the processing parameters, especially the sintering temperature. The objective of this study was to investigate the mechanical and thermal properties of pressureless-sintered SiC ceramics with a new quaternary sintering additive system (Al2O3Y2O3-CaO-SrO). The SiC ceramics were sintered at temperatures between 1750oC and 1950oC for 2 h in an argon atmosphere without an applied pressure. The present SiC ceramics showed process-tolerant behavior in terms of the microstructures and mechanical properties when the SiC ceramics were sintered at temperatures ranging from 1850oC to 1950oC. 2. Experimental procedure To prepare the SiC ceramics with Al2O3-Y2O3-CaO-SrO, 90.022 wt% β-SiC (~0.5 μm, grade BF-17, H.C. Starck, Berlin, Germany), 0.910 wt% α-SiC (~0.5 μm, FCP-15C, Norton AS, Lillesand, Norway), 5.042 wt% Al2O3 (~0.3 μm, AKP30, Sumitomo Chemical Co., Tokyo, Japan), 2.792 wt% Y2O3 (0.4 m, 99.99% pure, Kojundo Chemical Lab Co. Ltd., Sakado-shi, Japan), 0.433 wt% CaO, and 0.801 wt% SrO (98% pure, Kojundo Chemical Lab Co. Ltd.) were mixed by ball milling using SiC media in a polypropylene jar for 24 h in ethanol. The CaO was added in the form of CaCO3 (99.0% pure, Kojundo Chemical Lab Co. Ltd). The batch contained 93 vol% SiC and 7 vol% additives. The SiC powders consisted of 99 vol% βSiC matrix and 1 vol% α-SiC seed. The addition of the α-SiC seed into the β-SiC matrix usually accelerates the growth of platelet SiC grains [23,24]. The mixture was dried, sieved (60 mesh), uniaxially pressed into 40 × 40 × 5 mm samples at an applied pressure of 50 MPa, and subsequently cold pressed isostatically at 275 MPa. The samples were sintered at 17501950oC for 2 h in an argon atmosphere without any applied pressure. The heating rate was
20oC/min from 1000oC to the sintering temperature and the cooling rate was approximately 30oC/min. The relative density and apparent porosity of the sintered specimens were determined using the Archimedes method. The theoretical density of the specimen (3.289 g/cm3) was calculated according to the rule of mixtures. The sintered specimens were cut, polished, and etched with CF4 plasma containing 10% oxygen. The etched microstructure and fracture surface morphology were observed using scanning electron microscopy (SEM, S4300, Hitachi Ltd., Hitachi, Japan). The SEM micrographs of the specimens were quantitatively analyzed using an image analysis software package (Image-Pro Plus 4.0, Media Cybernetics Inc., MD, USA) using a procedure reported in a previous study [25]. The observed threedimensional morphology of the SiC grains was a hexagonal platelet in sintered specimens, where the intergranular phases were leached by a molten salt mixture. The definitions of the microstructural characteristics were described in a previous work [25]. The thickness of each grain (t) was determined directly from the shortest grain dimension in its two-dimensional image while the apparent length of each grain (L) was obtained from the largest dimension. The mean value of the highest 10% observed aspect ratios (L/t) was considered to be the mean of the actual values (R95) [26]. The grain thickness distribution was evaluated by plotting the areal frequency versus the equivalent grain thickness. The average thickness was taken as the value corresponding to one-half of the cumulative area. X-ray diffraction (XRD; D8 Discover, Bruker AXS GmbH, Karlsruhe, Germany) using Cu Kα radiation was performed on the starting SiC powder and ground powders of the sintered specimens. XRD data were analyzed using the Rietveld refinement method for quantitative phase analysis of the SiC polytypes.
The flexural strength was determined by ASTM C1161-13 [27]. The sintered samples were cut and ground into a 2.0 × 1.5 × 25 mm3 bar using an 800-grit diamond wheel for the flexural strength test. The tensile surfaces of the bars were polished to a 1-μm diamond finish and the tensile edges were chamfered to avoid stress concentration and large edge flaws caused by sectioning. Bending tests were performed at room temperature (RT) on six specimens for each condition using a four-point bending method with an inner and outer span of 10 and 20 mm, respectively, and a cross-head speed of 0.2 mm/min. The fracture toughness was determined according to ASTM C1421-16 [28]. The dimensions of the test specimens were 3 × 4 × 25 mm3. An indent was made in the middle of the polished surface of the test specimen using a Knoop indenter with a load of 49 N and a dwell time of 15 s. The residual stress damage zone was removed by mild polishing with 6 μm and subsequent 1 μm diamond slurries for 10 min each under a loading force of 0.2 kg before the fracture test. The Vickers hardness was measured using a Vickers indenter with a load of 4.9 N and a dwell time of 15 s. The thermal diffusivity was measured by the laser flash method. The details of the heat capacity (Cp) and the thermal diffusivity (α) measurements were described in a previous paper [29]. The α and Cp values were measured ten times each at RT and the average values were used to calculate the thermal conductivity, κ, using the following equation, κ = α ρ Cp
(1)
where ρ is the mass density of the specimen. The lattice oxygen content of the sintered specimens was measured using a hot-gas extraction method. The preparation of the specimens for the hot-gas extraction method was described in a previous work [21]. A commercial hot-gas extraction analyzer (model EMGA-920, Horiba Ltd., Kyoto, Japan) was used to determine the oxygen content. For comparison, the starting SiC powder was treated
using the same method as the sintered specimens before undergoing the same hot-gas extraction method. 3. Results and discussion 3. 1. Microstructural development and phase analysis The relative densities and apparent porosities of the SiC ceramics sintered with Al2O3-Y2O3CaO-SrO additives are shown in Table 1. The relative densities of the SiC ceramics were not very high due to significant evaporation of the sintering additives at high temperatures, leading to underestimation of the relative density of the specimens. The apparent porosity decreased from 7.0% to 0.3% with increasing sintering temperature from 1750oC to 1900oC, and then increased to 0.5% at 1950oC. Relative densities >95% can be achieved by sintering at temperatures between 1800oC and 1900oC for 2 h in an argon atmosphere without an applied pressure. The decrease of the sintered density at 1950oC was due to the increased weight loss by the partial evaporation of the liquid phase and the formation of volatile components such as Al2O and CO, which was caused by reactions between SiC and the sintering additives [30]. During sintering, the sintering additives (Al2O3-Y2O3-CaO-SrO) reacted with the native SiO2 film, which is always present on the surface of SiC particles, to form a Y-Sr-Ca-Al-Si-O melt. Subsequent heating at high temperatures led to the formation of a Y-Sr-Ca-Al-Si-O-C melt by the dissolution of SiC particles. The melt was responsible for the densification of the SiC ceramics via liquid phase sintering. The XRD patterns and quantitative phase analysis results of the SiC polytypes obtained using the Rietveld refinement method for the starting SiC powder and sintered specimens are shown in Fig. 1 and Table 2, respectively. The results show that (i) the significant β→α phase transformation of SiC occurred at 1750oC, (ii) the 3C→6H and the 6H→4H phase transformations gradually took place at temperatures ranging from 1750oC to 1850oC, (iii) all 3C-SiC transformed to 6H- and 4H-SiC at 1900oC, (iv) the 6H→4H phase transformation
took place at temperatures equal to or above 1900oC, and (v) the 3C-SiC content decreased with increasing sintering temperature, whereas the 4H-SiC content increased continuously with increasing sintering temperature. These results indicate that the 3C→6H→4H phase transformation of SiC took place with increasing sintering temperature. The present result is consistent with a previous work involving SiC ceramics sintered with Al2O3-Y2O3-CaO [21]. In the present work, appreciable amounts of the β→α phase transformation of SiC took place at a temperature as low as 1750oC, which is about 100-200oC lower than the specimens with binary Al2O3-Y2O3 additives [31,32]. A previous work involving SiC ceramics sintered with Al2O3-Y2O3-CaO [21] suggested that the melting point of the Al2O3-Y2O3-CaO-SiO2 system should be lower than that of the Al2O3-Y2O3-SiO2 system because the lowest eutectic temperature (1170oC) in the Al2O3-CaO-SiO2 system is lower than that (1350oC) in the Al2O3-Y2O3-SiO2 system. Thus, the melting point of the Al2O3-Y2O3-CaO-SrO-SiO2 system should also be lower than that of the Al2O3-Y2O3-SiO2 system. The lower melting point of the Al2O3-Y2O3-CaO-SrO-SiO2 system resulted in the acceleration of the β→α phase transformation of SiC at lower temperatures than SiC with Al2O3-Y2O3-CaO additives. The addition of the α-SiC seed into the β-SiC matrix is also known to accelerate the β→α phase transformation of SiC [20,25,33]. Therefore, the acceleration of the β→α phase transformation of SiC at a temperature as low as 1750oC was due to (i) the low viscosity of the Y-Sr-Ca-Si-Al-O-C melt due to its low melting point and (ii) the seeding effect of α-SiC added (1 vol%) in the present specimens. Typical microstructures of the SiC ceramics sintered with Al2O3-Y2O3-CaO-SrO are shown in Fig. 2. The 1800oC-sintered SiC ceramic consisted of equiaxed grains, whereas 1850- and 1950oC-sintered SiC ceramics consisted of relatively large elongated grains (hexagonal platelet grains in three dimensions) and small equiaxed grains. Toughened microstructures were achieved when the SiC ceramics were sintered at temperatures between 1850oC and
1950oC, by adding 1 vol% α-SiC seed into the β-SiC matrix. The present results are consistent with the microstructural development obtained in previous works [20,23,25] that contained α-SiC seeds. At ≥ 1900 °C the XRD-pattern and Table 2 show full transformation of β-SiC to the α-SiC. Since the microstructures contain small equiaxed grains, as shown in Fig. 2(c) and (d), the small equiaxed grains are considered as α-SiC, either 6H or 4H polytype. Keppeler et al. [34] and Ye et al. [35] also reported microstructures consisting of 100% α-SiC, containing small equiaxed grains. The microstructures also showed core/rim structures, indicating that the grain growth occurred through solution-reprecipitation [36,37]. It is well documented that both the core and the rim are composed of the same polytype [36]. Thus, α-phase was grown on the original α-SiC seeds. The rim area is the solid solution of SiC with soluble atoms in SiC lattice, i.e., Al and O in the present study. The average grain thickness increased from 1.0 μm to 1.9 μm with increasing sintering temperature from 1800oC to 1950oC. Figure 3 shows the grain thickness distribution of the SiC specimens. All specimens have a bimodal grain thickness distribution. The result is consistent with the microstructural observation (Fig. 2) which revealed relatively large elongated grains and small equiaxed grains. The bimodal grain thickness distribution was caused by the α-SiC seeding effect and low viscosity of the liquid phase formed from the current additive composition, which promoted the grain growth of platelet α-SiC grains during the liquid phase sintering. The increase of the sintering temperature moves the grain thickness distribution to a larger grain size region. The average aspect ratios of the SC80, SC85, SC90, and SC95 specimens were 2.6, 3.4, 3.8, and 3.8, respectively. The present result suggests that when Al2O3-Y2O3-CaO-SrO additives were used for densification of SiC, the aspect ratio of SiC grains in the resulting microstructure is insensitive to the sintering temperature change from 1850oC to 1950oC. The microstructures of the present SiC ceramics
were almost frozen during sintering at temperatures ranging from 1850oC to 1950oC. This phenomenon is similar to a previous work that investigated the process-tolerant behavior of SiC-TiC composites [38]. The frozen microstructure, i.e., sintering temperature-tolerant microstructure, was attributed to the impingement of elongated SiC grains and limited interaction of neighboring grains. The limited interaction of neighboring grains can be understood by the microstructural evolution principle suggested by Kang et al. [39,40]. According to the theory, the grain growth rate increases abruptly when the critical driving force for appreciable migration of a faceted boundary (Δgc) is almost equal to the maximum driving force for the largest grain (Δgmax). In contrast, stagnant grain growth is expected when Δgmax << Δgc. This means that almost the same microstructure can be obtained within a certain processing condition when Δgmax << Δgc. In the present work, when SiC ceramics were sintered at 1850-1950oC for 2 h in an argon atmosphere using Al2O3-Y2O3-CaO-SrO as sintering additives without an applied pressure, the Δgmax << Δgc condition was satisfied and stagnant grain growth took place, resulting in almost the same microstructure, i.e., sintering temperature-tolerant microstructure. 3. 2. Mechanical properties The variations of the fracture toughness and flexural strength of the SiC ceramics sintered with Al2O3-Y2O3-CaO-SrO as a function of the sintering temperature are shown in Fig. 4. The fracture toughness of the SiC ceramics increased from 4.1 to 4.9 MPa·m1/2 with increasing sintering temperature from 1800oC to 1850oC and reached a plateau in the temperature range of 1850oC to 1950oC. The higher fracture toughness of SC85 compared to SC80 is related to bridging and crack deflection by elongated SiC grains, as shown in Fig. 5. Generally, the mechanical properties of liquid-phase sintered SiC ceramics are strongly dependent on the additive composition, microstructure, and residual porosity [18,24,25,41].
Since the additive composition is the same for all specimens in this study, almost the same fracture toughness was obtained for the SC85-SC95 samples due to their similar microstructures. SC95 contained a smaller fraction of thicker grains than SC85 and SC90 (Fig. 3). Since thicker grains have a higher tendency for transgranular fracture than thinner grains due to their lower flexibility [25], the contribution of thicker grains to toughening was negligible, resulting in almost the same fracture toughness values of the SC85, SC90, and SC95 specimens. The fracture toughness values (4.9-5.0 MPa·m1/2) of the SC85, SC90, and SC95 specimens, which showed the toughened microstructure, are lower than the reported values (6-8 MPa·m1/2) for similar microstructures [20,24,25,42]. The difference is likely caused by the differences of the measurement methods. The toughness values reported in the previous studies were measured by the Vickers indentation method while the fracture toughness values obtained in the present work were determined by the surface crack in flexure method [28]. The fracture toughness values determined by the surface crack in flexure method are generally 1-2 MPa·m1/2 lower than those measured by the Vickers indentation method [21]. Figure 6 shows typical fracture surfaces of each specimen. All of the specimens exhibited mostly an intergranular fracture mode, which was the result of the weak interface created by the difference of the coefficients of thermal expansion of the liquid phase and SiC grains upon cooling after sintering. The fracture surfaces of SC85, SC90, and SC95 were more torturous than that of SC80. These results support the fracture toughness measurement results. The flexural strength values of SC80, SC85, SC90, and SC95 were 427, 454, 457, and 444 MPa, respectively. The fracture origins of the SiC ceramics were mostly surface pores or the agglomeration of several pores nearby surfaces, as shown in Fig. 7. The flexural strength values of pressureless-sintered SiC ceramics at room temperature reported in the literature include values of 341-556 MPa for SiC ceramics sintered with B-C [13,43,44], 321-532 MPa
for SiC ceramics sintered with Al2O3-Y2O3 [16,18,33,45,46], 418-555 MPa for SiC ceramics sintered with 10 vol% AlN-Yb2O3 [15], 437 MPa and 565 MPa for SiC ceramics sintered with 7 wt% Al2O3-CeO2 and Al2O3-Lu2O3, respectively [17], 399-433 MPa for SiC ceramics sintered with 3 vol% Al2O3-Y2O3-AlN [20], 377-440 MPa for SiC ceramics sintered with 10 wt% Al2O3-Y2O3-MgO [18], 516 MPa for SiC ceramics sintered with 12 wt% Al2O3-Y2O3TiO2 [19], 273-633 MPa for SiC ceramics sintered with 9 wt% Al2O3-Y2O3-CaO [21], and 509-520 MPa for SiC ceramics sintered with 6.5 vol% AlN-Y2O3-Sc2O3 [22]. The reported flexural strength values of pressureless-sintered SiC ceramics ranged from 273 MPa to 633 MPa, depending on the additive composition and processing conditions, i.e., microstructure. The flexural strength values (427-457 MPa) of the present ceramics are not excellent compared to previously reported values. However, the present ceramics sintered at temperatures ranging from 1850oC to 1950oC had almost the same fracture toughnesses (4.9, 5.0, and 4.9 MPa·m1/2 for SC85, SC90, and SC95, respectively) and flexural strength values (454, 457, and 444 MPa for SC85, SC90, and SC95, respectively) within the standard deviation, indicating insensitivity to the sintering temperature within the range of 18501950oC. As mentioned previously, the mechanical properties of SiC ceramics are strongly dependent on the additive composition, microstructure, and residual porosity [18,24,25,41]. The present SiC ceramics sintered at 1850-1950oC have almost the same apparent porosities (0.3-0.5%) with similar microstructures and aspect ratios (3.4-3.8). This small variation of the microstructure led to little variability of the mechanical properties (444-457 MPa and 4.9-5.0 MPa·m1/2) for the different sintering temperatures. The process tolerance, which is a measure of the variation of properties with respect to changes of the processing parameters, is a very important property for the mass production of ceramic parts. Generally, a temperature gradient is one of the main causes of a low manufacturing yield in mass production. A
temperature gradient in mass production generally leads to variability of the microstructure. This, in turn, leads to variability of the resulting properties, low manufacturing yields, and decreased reliability of the products. Thus, developing compositions that result in similar properties as the processing variables are varied, i.e., process-tolerant behavior, is very important for fabricating reliable products with high manufacturing yields. The present result suggests that the present sintering additive system (Al2O3-Y2O3-CaO-SrO), which exhibits process-tolerant behavior, should be beneficial for the mass production of SiC ceramic parts. 3.3. Thermal properties The thermal properties of the SiC ceramics as a function of sintering temperature including the thermal conductivity and the thermal diffusivity are shown in Fig. 8. The thermal diffusivity gradually increased from 27.9 to 37.1 mm2/s with increasing sintering temperature from 1800oC to 1900oC and then decreased slightly to 36.4 mm2/s when the SiC specimens were sintered at 1950oC. The thermal conductivity increased from 60.8 to 82.2 Wm-1K-1 with increasing sintering temperature from 1800oC to 1900oC. As the sintering temperature was further increased from 1900oC to 1950oC, the thermal conductivity decreased slightly to 77.1 Wm-1K-1. Generally, the heat is transported by phonons in SiC ceramics due to the lack of free electrons. The factors affecting the thermal conductivity of LPS-SiC ceramics include (1) the lattice oxygen content, which makes Si vacancies in the SiC lattice [47,48], (2) the amount of phase boundaries per unit volume, which is affected by the occurrence of the 3C→6H→4H phase transformation of SiC [21], (3) the number of grain boundaries per unit volume, which is related to the grain size, (4) the secondary phase in triple junctions, which has a much lower thermal conductivity than that of SiC grains [29], and (5) residual porosity, which results in enhanced phonon scattering at the pore-grain interfaces [49].
The increase of the thermal conductivity from 61 Wm-1K-1 to 82 Wm-1K-1 with increasing sintering temperature from 1800oC to 1900oC was due to (i) the decreased oxygen content in the SiC lattice from 0.366% to 0.253%, (ii) the decreased porosity from 0.8 to 0.3%, and (iii) the increased grain thickness from 1.0 μm to 1.4 μm. The thermal conductivity decreased from 82 Wm-1K-1 to 77 Wm-1K-1 with increasing sintering temperature from 1900oC to 1950oC, in spite of the decreased oxygen content in the SiC lattice from 0.253% to 0.247% and the increased grain thickness from 1.4 μm to 1.9 μm. The decrease of the thermal conductivity at 1950oC was due to the increased residual porosity from 0.3% to 0.5% and increased 6H→4H phase transformation of SiC [50]. The reported thermal conductivity of pressureless liquid phase-sintered SiC ceramics ranged from 55 to 110 Wm-1K-1, depending on the additive composition and microstructure: 55-70 Wm-1K-1 for SiC ceramics sintered with 10 vol% Al2O3-Y2O3 additives [51], 80 Wm-1K-1 for SiC ceramics sintered with 5 wt % Y3Al5O12-AlN additives [52], 68-73 Wm-1K-1 for SiC ceramics sintered with 10 wt % AlNY2O3 additives [53], 55-80 Wm-1K-1 for SiC ceramics sintered with 9 wt% Al2O3-Y2O3-CaO [21], and 92-110 Wm-1K-1 for SiC ceramics sintered with 6.5 vol% AlN-Y2O3-Sc2O3 additives [22]. The thermal conductivity of 82 Wm-1K-1 of the SC90 specimen is comparable to previously reported values and is lower than that of SiC ceramics sintered with 6.5 vol% AlN-Y2O3-Sc2O3 additives (92-110 Wm-1K-1). 4. Conclusions Dense SiC ceramics were fabricated using a new quaternary additive (Al2O3-Y2O3-CaO-SrO) by pressureless sintering at 1800-1950oC for 2 h in an argon atmosphere. Toughened microstructures could be achieved when the SiC ceramics were sintered at temperatures between 1850oC and 1950oC as a result of the beneficial effect of the present additive composition and the accelerated 3C→6H→4H phase transformation of SiC by 1 vol% α-SiC
seeding. The microstructures of the present SiC ceramics remained nearly the same within the temperature range of 1850-1950oC, due to the impingement of elongated SiC grains and limited interactions of neighboring grains. The flexural strength and fracture toughness of the present SiC ceramics were almost the same within the standard deviation when the SiC ceramics were sintered at temperatures between 1850oC and 1950oC (454 MPa and 4.9 MPa·m1/2 for SC85, 457 MPa and 5.0 MPa·m1/2 for SC90, and 444 MPa and 4.9 MPa·m1/2 for SC95, respectively). The small variability of the mechanical properties with respect to the sintering temperature, i.e., process-tolerant behavior, was attributed to the stagnant grain growth within the temperature range of 1850-1950oC. The present results suggest that SiC ceramics sintered with Al2O3Y2O3-CaO-SrO as a sintering additive are process-tolerant materials, especially insensitive to the sintering temperature. The improvement of the thermal conductivity from 61 Wm-1K-1 to 82 Wm-1K-1 with increasing sintering temperature from 1800oC to 1900oC was due to (i) the decreased lattice oxygen content from 0.366% to 0.253%, (ii) the decreased porosity from 0.8% to 0.3%, and (iii) the increased grain thickness from 1.0 to 1.4 μm at 1900oC. The typical flexural strength, fracture toughness, and thermal conductivity of the 1900oC-sintered SiC ceramics at RT were 457 MPa, 5.0 MPa∙m1/2, and 82 Wm-1K-1, respectively. Acknowledgements This work was supported by a grant from the National Research Foundation of Korea (NRF) funded by the Korea Government (MSIP) (2015R1A2A2A01004860). References [1] J. Rödel, A.B.N. Kounga, M. Weissenberger-Eibl, D. Koch, A. Bierwisch, W. Rossner, M.J. Hoffmann, R. Danzer, G. Schneider, Development of a roadmap for advanced ceramics: 2010–2025, J. Eur. Ceram. Soc. 29 (2009) 1549-1560.
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1750oC/2 h/Ar
Relative density (%) 92.8
Apparent porosity (%) 7.0
SC80
1800oC/2 h/Ar
95.2
0.8
SC85
1850oC/2 h/Ar
95.6
0.4
SC90
1900oC/2 h/Ar
96.7
0.3
SC95
1950oC/2 h/Ar
94.4
0.5
Sintering conditions
Table 2 Polytype contents in the Starting SiC powder and sintered SiC specimens sintered with Al2O3Y2O3-CaO-SrO. The polytype content was determined by Rietveld method. Polytype content (%) 3C 6H 4H 85.7 14.2 0.1
Sample designation
Sintering conditions
Starting powder
-
SC75
1750oC/2 h/Ar
60.3
22.9
18.8
SC80
1800oC/2 h/Ar
20.3
45.2
34.5
SC85
1850oC/2 h/Ar
5.7
45.0
49.3
SC90
1900oC/2 h/Ar
0.0
41.3
58.7
SC95
1950oC/2 h/Ar
0.0
39.0
61.0
Figure captions
Fig. 1. XRD patterns of the as-received SiC powder and SiC ceramics sintered with a new quaternary additive (Al2O3-Y2O3-CaO-SrO) at various temperatures without any applied pressure.
Fig. 2. Typical microstructures of SiC ceramics sintered at various temperatures for 2 h in argon: (a) 1800oC, (b) 1850oC, (c) 1900oC, and (d) 1950oC.
Fig. 3. Grain thickness distributions of SiC ceramics sintered at various temperatures for 2 h in argon: (a) 1800oC, (b) 1850oC, (c) 1900oC, and (d) 1950oC.
Fig. 4. Flexural strength and fracture toughness of SiC ceramics sintered with the Al2O3-Y2O3-CaOSrO additive as a function of the sintering temperature.
Fig. 5. SEM images of a crack path induced by a Vickers indenter for SiC ceramics sintered at various temperatures for 2 h in argon: (a) 1800 oC, (b) 1850oC, (c) 1900oC, and (d) 1950oC.
Fig. 6. Fracture surfaces of SiC ceramics sintered at various temperatures for 2 h in argon: (a) 1800oC, (b) 1850oC, (c) 1900oC, and (d) 1950oC.
Fig. 7. Typical fracture surfaces of SiC ceramics sintered at various temperatures for 2 h in argon: (a) 1800oC and (b) 1950oC.
Fig. 8. Thermal diffusivity and thermal conductivity of SiC ceramics sintered with a new quaternary additive (Al2O3-Y2O3-CaO-SrO) as a function of the sintering temperature.
S0955-2219(17)30599-X http://dx.doi.org/10.1016/j.jeurceramsoc.2017.09.011 JECS 11446
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Journal of the European Ceramic Society
Received date: Revised date: Accepted date:
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Please cite this article as: Seo Yu-Kwang, Eom Jung-Hye, Kim YoungWook.Process-tolerant pressureless-sintered silicon carbide ceramics with alumina-yttria-calcia-strontia.Journal of The European Ceramic Society http://dx.doi.org/10.1016/j.jeurceramsoc.2017.09.011 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Process-tolerant pressureless-sintered silicon carbide ceramics with aluminayttria-calcia-strontia Yu-Kwang Seo, Jung-Hye Eom, and Young-Wook Kim* Functional Ceramics Laboratory, Department of Materials Science and Engineering, The University of Seoul, Seoul 02504, Republic of Korea
Abstract Process-tolerant SiC ceramics were prepared by pressureless sintering at 1850-1950oC for 2 h in an argon atmosphere with a new quaternary additive (Al2O3-Y2O3-CaO-SrO). The SiC ceramics can be sintered to a >94% theoretical density at 1800-1950oC by pressureless sintering. Toughened microstructures consisting of relatively large platelet grains and small equiaxed grains were obtained when SiC ceramics were sintered at 1850-1950oC. The presently fabricated SiC ceramics showed little variability of the microstructure and mechanical properties with sintering within the temperature range of 1850-1950oC, demonstrating process-tolerant behavior. The thermal conductivity of the SiC ceramics increased with increasing sintering temperature from 1800oC to 1900oC due to decreases of the lattice oxygen content of the SiC grains and residual porosity. The flexural strength, fracture toughness, and thermal conductivity of the SiC ceramics sintered at 1850-1950oC were in the ranges of 444-457 MPa, 4.9-5.0 MPa∙m1/2, and 76-82 Wm-1K-1, respectively. Keywords: Silicon carbide; Mechanical properties; Thermal conductivity; Microstructure; Process-tolerant behavior *Corresponding author. Fax: +82 2 6490 2404 E-mail address: [email protected] (Y.-W. Kim). 1. Introduction The market for advanced ceramics is continuously growing due to their excellent performance, multifunctional properties, and high energy efficiency [1]. However, the lower yields for ceramics than for those of other materials such as metals and polymers could be a
barrier for widening the industrial applications of advanced ceramics due to cost effectiveness [2,3]. Processing parameters such as the sintering temperature, time, and annealing conditions affect the microstructural characteristics and consequently lead to variability of the electrical, thermal, optical and mechanical properties [4-8]. Variability of these properties can potentially lower the production yield. Thus, finding systems that exhibit “process-tolerant behavior,” i.e., the insensitivity of properties to specific processing parameters, is an important issue in ceramic industries for improving the production yield and cost effectiveness [9]. Silicon carbide (SiC) is one of the most important engineering ceramics due to its good thermal conductivity, good oxidation and corrosion resistance, as well as its excellent mechanical properties [5,7,10-12]. Thus, dense SiC ceramics are widely used in many applications such as heaters, heat exchangers, semiconductor processing parts, mechanical seals, and cladding materials for nuclear fuels. However, the densification of SiC ceramics without sintering additives is difficult due to their highly covalent bonding and low selfdiffusivity. Various sintering additive compositions were investigated for pressurelesssintered SiC ceramics. The relative density and flexural strength of pressureless-sintered SiC ceramics showed wide variations (96–99% and 394–633 MPa, respectively), which depended on the additive composition and processing conditions [13-22]. Previous results include values of 98% and 556 MPa for SiC ceramics sintered with 3.5 wt% B-C [13], 99% and 394 MPa for SiC ceramics sintered with 7 wt% AlN-C [14], 96% and 440 MPa for SiC ceramics sintered with 10 wt% Al2O3-Y2O3-MgO [18], 98% and 433 MPa for SiC ceramics sintered with 3 vol% Al2O3-AlN-Y2O3 [20], 98% and 633 MPa for SiC ceramics sintered with 9 wt% Al2O3-Y2O3-CaO [21], and 97% and 520 MPa for SiC ceramics sintered with 6.5 vol% AlNY2O3-Sc2O3 [22]. However, there is no previous work investigating pressureless-sintered SiC ceramics with a quaternary additive system. The reported pressureless-sintered SiC ceramics
with binary or ternary sintering additives showed a high sintered density and excellent mechanical properties. However, all of the above pressureless-sintered SiC ceramics showed a wide variation of microstructures and mechanical properties depending on the processing parameters, especially the sintering temperature. The objective of this study was to investigate the mechanical and thermal properties of pressureless-sintered SiC ceramics with a new quaternary sintering additive system (Al2O3Y2O3-CaO-SrO). The SiC ceramics were sintered at temperatures between 1750oC and 1950oC for 2 h in an argon atmosphere without an applied pressure. The present SiC ceramics showed process-tolerant behavior in terms of the microstructures and mechanical properties when the SiC ceramics were sintered at temperatures ranging from 1850oC to 1950oC. 2. Experimental procedure To prepare the SiC ceramics with Al2O3-Y2O3-CaO-SrO, 90.022 wt% β-SiC (~0.5 μm, grade BF-17, H.C. Starck, Berlin, Germany), 0.910 wt% α-SiC (~0.5 μm, FCP-15C, Norton AS, Lillesand, Norway), 5.042 wt% Al2O3 (~0.3 μm, AKP30, Sumitomo Chemical Co., Tokyo, Japan), 2.792 wt% Y2O3 (0.4 m, 99.99% pure, Kojundo Chemical Lab Co. Ltd., Sakado-shi, Japan), 0.433 wt% CaO, and 0.801 wt% SrO (98% pure, Kojundo Chemical Lab Co. Ltd.) were mixed by ball milling using SiC media in a polypropylene jar for 24 h in ethanol. The CaO was added in the form of CaCO3 (99.0% pure, Kojundo Chemical Lab Co. Ltd). The batch contained 93 vol% SiC and 7 vol% additives. The SiC powders consisted of 99 vol% βSiC matrix and 1 vol% α-SiC seed. The addition of the α-SiC seed into the β-SiC matrix usually accelerates the growth of platelet SiC grains [23,24]. The mixture was dried, sieved (60 mesh), uniaxially pressed into 40 × 40 × 5 mm samples at an applied pressure of 50 MPa, and subsequently cold pressed isostatically at 275 MPa. The samples were sintered at 17501950oC for 2 h in an argon atmosphere without any applied pressure. The heating rate was
20oC/min from 1000oC to the sintering temperature and the cooling rate was approximately 30oC/min. The relative density and apparent porosity of the sintered specimens were determined using the Archimedes method. The theoretical density of the specimen (3.289 g/cm3) was calculated according to the rule of mixtures. The sintered specimens were cut, polished, and etched with CF4 plasma containing 10% oxygen. The etched microstructure and fracture surface morphology were observed using scanning electron microscopy (SEM, S4300, Hitachi Ltd., Hitachi, Japan). The SEM micrographs of the specimens were quantitatively analyzed using an image analysis software package (Image-Pro Plus 4.0, Media Cybernetics Inc., MD, USA) using a procedure reported in a previous study [25]. The observed threedimensional morphology of the SiC grains was a hexagonal platelet in sintered specimens, where the intergranular phases were leached by a molten salt mixture. The definitions of the microstructural characteristics were described in a previous work [25]. The thickness of each grain (t) was determined directly from the shortest grain dimension in its two-dimensional image while the apparent length of each grain (L) was obtained from the largest dimension. The mean value of the highest 10% observed aspect ratios (L/t) was considered to be the mean of the actual values (R95) [26]. The grain thickness distribution was evaluated by plotting the areal frequency versus the equivalent grain thickness. The average thickness was taken as the value corresponding to one-half of the cumulative area. X-ray diffraction (XRD; D8 Discover, Bruker AXS GmbH, Karlsruhe, Germany) using Cu Kα radiation was performed on the starting SiC powder and ground powders of the sintered specimens. XRD data were analyzed using the Rietveld refinement method for quantitative phase analysis of the SiC polytypes.
The flexural strength was determined by ASTM C1161-13 [27]. The sintered samples were cut and ground into a 2.0 × 1.5 × 25 mm3 bar using an 800-grit diamond wheel for the flexural strength test. The tensile surfaces of the bars were polished to a 1-μm diamond finish and the tensile edges were chamfered to avoid stress concentration and large edge flaws caused by sectioning. Bending tests were performed at room temperature (RT) on six specimens for each condition using a four-point bending method with an inner and outer span of 10 and 20 mm, respectively, and a cross-head speed of 0.2 mm/min. The fracture toughness was determined according to ASTM C1421-16 [28]. The dimensions of the test specimens were 3 × 4 × 25 mm3. An indent was made in the middle of the polished surface of the test specimen using a Knoop indenter with a load of 49 N and a dwell time of 15 s. The residual stress damage zone was removed by mild polishing with 6 μm and subsequent 1 μm diamond slurries for 10 min each under a loading force of 0.2 kg before the fracture test. The Vickers hardness was measured using a Vickers indenter with a load of 4.9 N and a dwell time of 15 s. The thermal diffusivity was measured by the laser flash method. The details of the heat capacity (Cp) and the thermal diffusivity (α) measurements were described in a previous paper [29]. The α and Cp values were measured ten times each at RT and the average values were used to calculate the thermal conductivity, κ, using the following equation, κ = α ρ Cp
(1)
where ρ is the mass density of the specimen. The lattice oxygen content of the sintered specimens was measured using a hot-gas extraction method. The preparation of the specimens for the hot-gas extraction method was described in a previous work [21]. A commercial hot-gas extraction analyzer (model EMGA-920, Horiba Ltd., Kyoto, Japan) was used to determine the oxygen content. For comparison, the starting SiC powder was treated
using the same method as the sintered specimens before undergoing the same hot-gas extraction method. 3. Results and discussion 3. 1. Microstructural development and phase analysis The relative densities and apparent porosities of the SiC ceramics sintered with Al2O3-Y2O3CaO-SrO additives are shown in Table 1. The relative densities of the SiC ceramics were not very high due to significant evaporation of the sintering additives at high temperatures, leading to underestimation of the relative density of the specimens. The apparent porosity decreased from 7.0% to 0.3% with increasing sintering temperature from 1750oC to 1900oC, and then increased to 0.5% at 1950oC. Relative densities >95% can be achieved by sintering at temperatures between 1800oC and 1900oC for 2 h in an argon atmosphere without an applied pressure. The decrease of the sintered density at 1950oC was due to the increased weight loss by the partial evaporation of the liquid phase and the formation of volatile components such as Al2O and CO, which was caused by reactions between SiC and the sintering additives [30]. During sintering, the sintering additives (Al2O3-Y2O3-CaO-SrO) reacted with the native SiO2 film, which is always present on the surface of SiC particles, to form a Y-Sr-Ca-Al-Si-O melt. Subsequent heating at high temperatures led to the formation of a Y-Sr-Ca-Al-Si-O-C melt by the dissolution of SiC particles. The melt was responsible for the densification of the SiC ceramics via liquid phase sintering. The XRD patterns and quantitative phase analysis results of the SiC polytypes obtained using the Rietveld refinement method for the starting SiC powder and sintered specimens are shown in Fig. 1 and Table 2, respectively. The results show that (i) the significant β→α phase transformation of SiC occurred at 1750oC, (ii) the 3C→6H and the 6H→4H phase transformations gradually took place at temperatures ranging from 1750oC to 1850oC, (iii) all 3C-SiC transformed to 6H- and 4H-SiC at 1900oC, (iv) the 6H→4H phase transformation
took place at temperatures equal to or above 1900oC, and (v) the 3C-SiC content decreased with increasing sintering temperature, whereas the 4H-SiC content increased continuously with increasing sintering temperature. These results indicate that the 3C→6H→4H phase transformation of SiC took place with increasing sintering temperature. The present result is consistent with a previous work involving SiC ceramics sintered with Al2O3-Y2O3-CaO [21]. In the present work, appreciable amounts of the β→α phase transformation of SiC took place at a temperature as low as 1750oC, which is about 100-200oC lower than the specimens with binary Al2O3-Y2O3 additives [31,32]. A previous work involving SiC ceramics sintered with Al2O3-Y2O3-CaO [21] suggested that the melting point of the Al2O3-Y2O3-CaO-SiO2 system should be lower than that of the Al2O3-Y2O3-SiO2 system because the lowest eutectic temperature (1170oC) in the Al2O3-CaO-SiO2 system is lower than that (1350oC) in the Al2O3-Y2O3-SiO2 system. Thus, the melting point of the Al2O3-Y2O3-CaO-SrO-SiO2 system should also be lower than that of the Al2O3-Y2O3-SiO2 system. The lower melting point of the Al2O3-Y2O3-CaO-SrO-SiO2 system resulted in the acceleration of the β→α phase transformation of SiC at lower temperatures than SiC with Al2O3-Y2O3-CaO additives. The addition of the α-SiC seed into the β-SiC matrix is also known to accelerate the β→α phase transformation of SiC [20,25,33]. Therefore, the acceleration of the β→α phase transformation of SiC at a temperature as low as 1750oC was due to (i) the low viscosity of the Y-Sr-Ca-Si-Al-O-C melt due to its low melting point and (ii) the seeding effect of α-SiC added (1 vol%) in the present specimens. Typical microstructures of the SiC ceramics sintered with Al2O3-Y2O3-CaO-SrO are shown in Fig. 2. The 1800oC-sintered SiC ceramic consisted of equiaxed grains, whereas 1850- and 1950oC-sintered SiC ceramics consisted of relatively large elongated grains (hexagonal platelet grains in three dimensions) and small equiaxed grains. Toughened microstructures were achieved when the SiC ceramics were sintered at temperatures between 1850oC and
1950oC, by adding 1 vol% α-SiC seed into the β-SiC matrix. The present results are consistent with the microstructural development obtained in previous works [20,23,25] that contained α-SiC seeds. At ≥ 1900 °C the XRD-pattern and Table 2 show full transformation of β-SiC to the α-SiC. Since the microstructures contain small equiaxed grains, as shown in Fig. 2(c) and (d), the small equiaxed grains are considered as α-SiC, either 6H or 4H polytype. Keppeler et al. [34] and Ye et al. [35] also reported microstructures consisting of 100% α-SiC, containing small equiaxed grains. The microstructures also showed core/rim structures, indicating that the grain growth occurred through solution-reprecipitation [36,37]. It is well documented that both the core and the rim are composed of the same polytype [36]. Thus, α-phase was grown on the original α-SiC seeds. The rim area is the solid solution of SiC with soluble atoms in SiC lattice, i.e., Al and O in the present study. The average grain thickness increased from 1.0 μm to 1.9 μm with increasing sintering temperature from 1800oC to 1950oC. Figure 3 shows the grain thickness distribution of the SiC specimens. All specimens have a bimodal grain thickness distribution. The result is consistent with the microstructural observation (Fig. 2) which revealed relatively large elongated grains and small equiaxed grains. The bimodal grain thickness distribution was caused by the α-SiC seeding effect and low viscosity of the liquid phase formed from the current additive composition, which promoted the grain growth of platelet α-SiC grains during the liquid phase sintering. The increase of the sintering temperature moves the grain thickness distribution to a larger grain size region. The average aspect ratios of the SC80, SC85, SC90, and SC95 specimens were 2.6, 3.4, 3.8, and 3.8, respectively. The present result suggests that when Al2O3-Y2O3-CaO-SrO additives were used for densification of SiC, the aspect ratio of SiC grains in the resulting microstructure is insensitive to the sintering temperature change from 1850oC to 1950oC. The microstructures of the present SiC ceramics
were almost frozen during sintering at temperatures ranging from 1850oC to 1950oC. This phenomenon is similar to a previous work that investigated the process-tolerant behavior of SiC-TiC composites [38]. The frozen microstructure, i.e., sintering temperature-tolerant microstructure, was attributed to the impingement of elongated SiC grains and limited interaction of neighboring grains. The limited interaction of neighboring grains can be understood by the microstructural evolution principle suggested by Kang et al. [39,40]. According to the theory, the grain growth rate increases abruptly when the critical driving force for appreciable migration of a faceted boundary (Δgc) is almost equal to the maximum driving force for the largest grain (Δgmax). In contrast, stagnant grain growth is expected when Δgmax << Δgc. This means that almost the same microstructure can be obtained within a certain processing condition when Δgmax << Δgc. In the present work, when SiC ceramics were sintered at 1850-1950oC for 2 h in an argon atmosphere using Al2O3-Y2O3-CaO-SrO as sintering additives without an applied pressure, the Δgmax << Δgc condition was satisfied and stagnant grain growth took place, resulting in almost the same microstructure, i.e., sintering temperature-tolerant microstructure. 3. 2. Mechanical properties The variations of the fracture toughness and flexural strength of the SiC ceramics sintered with Al2O3-Y2O3-CaO-SrO as a function of the sintering temperature are shown in Fig. 4. The fracture toughness of the SiC ceramics increased from 4.1 to 4.9 MPa·m1/2 with increasing sintering temperature from 1800oC to 1850oC and reached a plateau in the temperature range of 1850oC to 1950oC. The higher fracture toughness of SC85 compared to SC80 is related to bridging and crack deflection by elongated SiC grains, as shown in Fig. 5. Generally, the mechanical properties of liquid-phase sintered SiC ceramics are strongly dependent on the additive composition, microstructure, and residual porosity [18,24,25,41].
Since the additive composition is the same for all specimens in this study, almost the same fracture toughness was obtained for the SC85-SC95 samples due to their similar microstructures. SC95 contained a smaller fraction of thicker grains than SC85 and SC90 (Fig. 3). Since thicker grains have a higher tendency for transgranular fracture than thinner grains due to their lower flexibility [25], the contribution of thicker grains to toughening was negligible, resulting in almost the same fracture toughness values of the SC85, SC90, and SC95 specimens. The fracture toughness values (4.9-5.0 MPa·m1/2) of the SC85, SC90, and SC95 specimens, which showed the toughened microstructure, are lower than the reported values (6-8 MPa·m1/2) for similar microstructures [20,24,25,42]. The difference is likely caused by the differences of the measurement methods. The toughness values reported in the previous studies were measured by the Vickers indentation method while the fracture toughness values obtained in the present work were determined by the surface crack in flexure method [28]. The fracture toughness values determined by the surface crack in flexure method are generally 1-2 MPa·m1/2 lower than those measured by the Vickers indentation method [21]. Figure 6 shows typical fracture surfaces of each specimen. All of the specimens exhibited mostly an intergranular fracture mode, which was the result of the weak interface created by the difference of the coefficients of thermal expansion of the liquid phase and SiC grains upon cooling after sintering. The fracture surfaces of SC85, SC90, and SC95 were more torturous than that of SC80. These results support the fracture toughness measurement results. The flexural strength values of SC80, SC85, SC90, and SC95 were 427, 454, 457, and 444 MPa, respectively. The fracture origins of the SiC ceramics were mostly surface pores or the agglomeration of several pores nearby surfaces, as shown in Fig. 7. The flexural strength values of pressureless-sintered SiC ceramics at room temperature reported in the literature include values of 341-556 MPa for SiC ceramics sintered with B-C [13,43,44], 321-532 MPa
for SiC ceramics sintered with Al2O3-Y2O3 [16,18,33,45,46], 418-555 MPa for SiC ceramics sintered with 10 vol% AlN-Yb2O3 [15], 437 MPa and 565 MPa for SiC ceramics sintered with 7 wt% Al2O3-CeO2 and Al2O3-Lu2O3, respectively [17], 399-433 MPa for SiC ceramics sintered with 3 vol% Al2O3-Y2O3-AlN [20], 377-440 MPa for SiC ceramics sintered with 10 wt% Al2O3-Y2O3-MgO [18], 516 MPa for SiC ceramics sintered with 12 wt% Al2O3-Y2O3TiO2 [19], 273-633 MPa for SiC ceramics sintered with 9 wt% Al2O3-Y2O3-CaO [21], and 509-520 MPa for SiC ceramics sintered with 6.5 vol% AlN-Y2O3-Sc2O3 [22]. The reported flexural strength values of pressureless-sintered SiC ceramics ranged from 273 MPa to 633 MPa, depending on the additive composition and processing conditions, i.e., microstructure. The flexural strength values (427-457 MPa) of the present ceramics are not excellent compared to previously reported values. However, the present ceramics sintered at temperatures ranging from 1850oC to 1950oC had almost the same fracture toughnesses (4.9, 5.0, and 4.9 MPa·m1/2 for SC85, SC90, and SC95, respectively) and flexural strength values (454, 457, and 444 MPa for SC85, SC90, and SC95, respectively) within the standard deviation, indicating insensitivity to the sintering temperature within the range of 18501950oC. As mentioned previously, the mechanical properties of SiC ceramics are strongly dependent on the additive composition, microstructure, and residual porosity [18,24,25,41]. The present SiC ceramics sintered at 1850-1950oC have almost the same apparent porosities (0.3-0.5%) with similar microstructures and aspect ratios (3.4-3.8). This small variation of the microstructure led to little variability of the mechanical properties (444-457 MPa and 4.9-5.0 MPa·m1/2) for the different sintering temperatures. The process tolerance, which is a measure of the variation of properties with respect to changes of the processing parameters, is a very important property for the mass production of ceramic parts. Generally, a temperature gradient is one of the main causes of a low manufacturing yield in mass production. A
temperature gradient in mass production generally leads to variability of the microstructure. This, in turn, leads to variability of the resulting properties, low manufacturing yields, and decreased reliability of the products. Thus, developing compositions that result in similar properties as the processing variables are varied, i.e., process-tolerant behavior, is very important for fabricating reliable products with high manufacturing yields. The present result suggests that the present sintering additive system (Al2O3-Y2O3-CaO-SrO), which exhibits process-tolerant behavior, should be beneficial for the mass production of SiC ceramic parts. 3.3. Thermal properties The thermal properties of the SiC ceramics as a function of sintering temperature including the thermal conductivity and the thermal diffusivity are shown in Fig. 8. The thermal diffusivity gradually increased from 27.9 to 37.1 mm2/s with increasing sintering temperature from 1800oC to 1900oC and then decreased slightly to 36.4 mm2/s when the SiC specimens were sintered at 1950oC. The thermal conductivity increased from 60.8 to 82.2 Wm-1K-1 with increasing sintering temperature from 1800oC to 1900oC. As the sintering temperature was further increased from 1900oC to 1950oC, the thermal conductivity decreased slightly to 77.1 Wm-1K-1. Generally, the heat is transported by phonons in SiC ceramics due to the lack of free electrons. The factors affecting the thermal conductivity of LPS-SiC ceramics include (1) the lattice oxygen content, which makes Si vacancies in the SiC lattice [47,48], (2) the amount of phase boundaries per unit volume, which is affected by the occurrence of the 3C→6H→4H phase transformation of SiC [21], (3) the number of grain boundaries per unit volume, which is related to the grain size, (4) the secondary phase in triple junctions, which has a much lower thermal conductivity than that of SiC grains [29], and (5) residual porosity, which results in enhanced phonon scattering at the pore-grain interfaces [49].
The increase of the thermal conductivity from 61 Wm-1K-1 to 82 Wm-1K-1 with increasing sintering temperature from 1800oC to 1900oC was due to (i) the decreased oxygen content in the SiC lattice from 0.366% to 0.253%, (ii) the decreased porosity from 0.8 to 0.3%, and (iii) the increased grain thickness from 1.0 μm to 1.4 μm. The thermal conductivity decreased from 82 Wm-1K-1 to 77 Wm-1K-1 with increasing sintering temperature from 1900oC to 1950oC, in spite of the decreased oxygen content in the SiC lattice from 0.253% to 0.247% and the increased grain thickness from 1.4 μm to 1.9 μm. The decrease of the thermal conductivity at 1950oC was due to the increased residual porosity from 0.3% to 0.5% and increased 6H→4H phase transformation of SiC [50]. The reported thermal conductivity of pressureless liquid phase-sintered SiC ceramics ranged from 55 to 110 Wm-1K-1, depending on the additive composition and microstructure: 55-70 Wm-1K-1 for SiC ceramics sintered with 10 vol% Al2O3-Y2O3 additives [51], 80 Wm-1K-1 for SiC ceramics sintered with 5 wt % Y3Al5O12-AlN additives [52], 68-73 Wm-1K-1 for SiC ceramics sintered with 10 wt % AlNY2O3 additives [53], 55-80 Wm-1K-1 for SiC ceramics sintered with 9 wt% Al2O3-Y2O3-CaO [21], and 92-110 Wm-1K-1 for SiC ceramics sintered with 6.5 vol% AlN-Y2O3-Sc2O3 additives [22]. The thermal conductivity of 82 Wm-1K-1 of the SC90 specimen is comparable to previously reported values and is lower than that of SiC ceramics sintered with 6.5 vol% AlN-Y2O3-Sc2O3 additives (92-110 Wm-1K-1). 4. Conclusions Dense SiC ceramics were fabricated using a new quaternary additive (Al2O3-Y2O3-CaO-SrO) by pressureless sintering at 1800-1950oC for 2 h in an argon atmosphere. Toughened microstructures could be achieved when the SiC ceramics were sintered at temperatures between 1850oC and 1950oC as a result of the beneficial effect of the present additive composition and the accelerated 3C→6H→4H phase transformation of SiC by 1 vol% α-SiC
seeding. The microstructures of the present SiC ceramics remained nearly the same within the temperature range of 1850-1950oC, due to the impingement of elongated SiC grains and limited interactions of neighboring grains. The flexural strength and fracture toughness of the present SiC ceramics were almost the same within the standard deviation when the SiC ceramics were sintered at temperatures between 1850oC and 1950oC (454 MPa and 4.9 MPa·m1/2 for SC85, 457 MPa and 5.0 MPa·m1/2 for SC90, and 444 MPa and 4.9 MPa·m1/2 for SC95, respectively). The small variability of the mechanical properties with respect to the sintering temperature, i.e., process-tolerant behavior, was attributed to the stagnant grain growth within the temperature range of 1850-1950oC. The present results suggest that SiC ceramics sintered with Al2O3Y2O3-CaO-SrO as a sintering additive are process-tolerant materials, especially insensitive to the sintering temperature. The improvement of the thermal conductivity from 61 Wm-1K-1 to 82 Wm-1K-1 with increasing sintering temperature from 1800oC to 1900oC was due to (i) the decreased lattice oxygen content from 0.366% to 0.253%, (ii) the decreased porosity from 0.8% to 0.3%, and (iii) the increased grain thickness from 1.0 to 1.4 μm at 1900oC. The typical flexural strength, fracture toughness, and thermal conductivity of the 1900oC-sintered SiC ceramics at RT were 457 MPa, 5.0 MPa∙m1/2, and 82 Wm-1K-1, respectively. Acknowledgements This work was supported by a grant from the National Research Foundation of Korea (NRF) funded by the Korea Government (MSIP) (2015R1A2A2A01004860). References [1] J. Rödel, A.B.N. Kounga, M. Weissenberger-Eibl, D. Koch, A. Bierwisch, W. Rossner, M.J. Hoffmann, R. Danzer, G. Schneider, Development of a roadmap for advanced ceramics: 2010–2025, J. Eur. Ceram. Soc. 29 (2009) 1549-1560.
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[50] Y. Zhou, K. Hirao, Y. Yamauchi, S. Kanzaki, Effects of rare-earth oxide and alumina ad ditives on thermal conductivity of liquid-phase-sintered silicon carbide, J. Mater. Res. 1 8 (2003) 1854-1862. [51] E. Volz, A. Roosen, W. Hartung, A. Winnacker, Electrical and thermal conductivity of li quid phase sintered SiC, J. Eur. Ceram. Soc. 21 (2001) 2089-2093. [52] L.S. Sigl, Thermal conductivity of liquid phase sintered silicon carbide, J. Eur. Ceram. S oc. 23 (2003) 1115-1122. [53] V.P. Onbattuvelli, R.K. Enneti, S.V. Atre, The effects of nanoparticle addition on the den sification and properties of SiC, Ceram. Int. 38 (2012) 5393-5399. Table 1 Relative density and apparent porosity of SiC ceramics sintered with Al2O3-Y2O3-CaO-SrO. Sample designation SC75
1750oC/2 h/Ar
Relative density (%) 92.8
Apparent porosity (%) 7.0
SC80
1800oC/2 h/Ar
95.2
0.8
SC85
1850oC/2 h/Ar
95.6
0.4
SC90
1900oC/2 h/Ar
96.7
0.3
SC95
1950oC/2 h/Ar
94.4
0.5
Sintering conditions
Table 2 Polytype contents in the Starting SiC powder and sintered SiC specimens sintered with Al2O3Y2O3-CaO-SrO. The polytype content was determined by Rietveld method. Polytype content (%) 3C 6H 4H 85.7 14.2 0.1
Sample designation
Sintering conditions
Starting powder
-
SC75
1750oC/2 h/Ar
60.3
22.9
18.8
SC80
1800oC/2 h/Ar
20.3
45.2
34.5
SC85
1850oC/2 h/Ar
5.7
45.0
49.3
SC90
1900oC/2 h/Ar
0.0
41.3
58.7
SC95
1950oC/2 h/Ar
0.0
39.0
61.0
Figure captions
Fig. 1. XRD patterns of the as-received SiC powder and SiC ceramics sintered with a new quaternary additive (Al2O3-Y2O3-CaO-SrO) at various temperatures without any applied pressure.
Fig. 2. Typical microstructures of SiC ceramics sintered at various temperatures for 2 h in argon: (a) 1800oC, (b) 1850oC, (c) 1900oC, and (d) 1950oC.
Fig. 3. Grain thickness distributions of SiC ceramics sintered at various temperatures for 2 h in argon: (a) 1800oC, (b) 1850oC, (c) 1900oC, and (d) 1950oC.
Fig. 4. Flexural strength and fracture toughness of SiC ceramics sintered with the Al2O3-Y2O3-CaOSrO additive as a function of the sintering temperature.
Fig. 5. SEM images of a crack path induced by a Vickers indenter for SiC ceramics sintered at various temperatures for 2 h in argon: (a) 1800 oC, (b) 1850oC, (c) 1900oC, and (d) 1950oC.
Fig. 6. Fracture surfaces of SiC ceramics sintered at various temperatures for 2 h in argon: (a) 1800oC, (b) 1850oC, (c) 1900oC, and (d) 1950oC.
Fig. 7. Typical fracture surfaces of SiC ceramics sintered at various temperatures for 2 h in argon: (a) 1800oC and (b) 1950oC.
Fig. 8. Thermal diffusivity and thermal conductivity of SiC ceramics sintered with a new quaternary additive (Al2O3-Y2O3-CaO-SrO) as a function of the sintering temperature.