Abstract: The SiC/Al4SiC4 composites with the improved mechanical properties and thermal conductivity were fabricated by the in-situ reaction of polycarbosilane (PCS) and Al powders using spark plasma sintering. The addition of 5 wt% yttrium (Y) sintering additive was useful to obtain fully dense samples after sintering at a relatively low temperature of 1650 ℃, due to the formation of a liquid phase during sintering. The average particle size of the in-situ formed SiC was ~300 nm. The fracture toughness (4.9 MPa·m1/2), Vickers hardness (16.3 GPa), and thermal conductivity (15.8 W/(m·K)) of the SiC/Al4SiC4 composite sintered at 1650 ℃ were significantly higher than the hardness (13.2 GPa), fracture toughness (2.16 MPa·m1/2), and thermal conductivity (7.8 W/(m·K)) of the monolithic Al4SiC4 ceramics. The improved mechanical and thermal properties of the composites were attributed to the high density, fine grain size, as well as the optimized grain boundary structure of the SiC/Al4SiC4 composites.
Keywords: Al4SiC4; SiC; Y3Si2C2; spark plasma sintering
Aluminum silicon carbide (Al4SiC4) is ternary layered ceramics with hexagonal structure symmetry of space group P63mc (No. 186), in which the Al4C3-type and SiC-type sub-units are alternatively stacked along the c-axis direction . Al4SiC4 has good combination of properties, such as low density (3.03 g/cm³), high melting point (2080 ℃), high compressive strength (260 GPa), as well as remarkable resistance to oxidation, corrosion, and hydration at high temperatures [2–6]. Due to its good oxidation resistance in the saturated water vapor environments [7,8], high strength at elevated temperatures , and excellent slag corrosion resistance [10,11], Al4SiC4 has potential to be used in aerospace and nuclear applications. However, its relatively low hardness (10.6 GPa) and fracture toughness (3.98 MPa·m1/2)  may limit its application .
One efficient way to improve hardness, fracture toughness, and thermal conductivity of Al4SiC4 is to incorporate a reinforcing filler with high thermal conductivity into the Al4SiC4 matrix. Silicon carbide (SiC) is a promising candidate, because of its high hardness, excellent mechanical properties at high temperatures, and high thermal conductivity [12–15]. The SiC particles have successfully been used to reinforce titanium silicon carbide (Ti3SiC2) matrices [16–18]. Zhang et al.  reported that the hardness, fracture toughness, and thermal conductivity of Ti3SiC2 were improved by the addition of SiC nano particles into the Ti3SiC2 matrix. Therefore, the motivation of the present work was to utilize the reinforcing effect of SiC nano particles to develop the SiC/Al4SiC4 composites with improved mechanical properties and thermal conductivity.
Pure Al4SiC4 can be synthesized using different methods and starting materials, which mainly include: 1) solid-state reactions of the starting mixture containing aluminum, silicon, and carbon powders [6,19,20]; 2) solid-state reactions of the aluminum, carbon, and polycarbosilane (PCS) starting mixture, in which PCS replaces Si as the source of Si and C ; 3) solid-state reactions of the starting Al4C3 and SiC powders [21,22]; 4) carbothermal reduction of the Al2O3 or Al(OH)3, SiO2, and C powder mixture [23,24].
Among all of these methods, the one that uses PCS as a raw material attracted our attention to develop SiC/Al4SiC4 composites. This is because the nano sized SiC particles, which are beneficial for toughening the Al4SiC4 materials, could be derived from PCS. On the other hand, the densification of Al4SiC4 is usually performed at high temperatures, around 1900 ℃ . The use of sintering additives is necessary to reduce the sintering temperature. Recently, yttrium silicon carbide (Y3Si2C2) was successfully used as the sintering additive for SiC  and predicted as a promising interphase material with great interest for SiCf/SiC composites . According to the calculated Y–Si–C ternary phase diagram, a liquid phase is formed at the temperatures of ~1560 ℃ . The presence of the liquid phase effectively promotes densification of SiC ceramics. Furthermore, an eutectic liquid phase can also be formed in the Y2O3–Al2O3–SiO2 system at the temperature of ~1640 ℃ . From this point of view, two different liquid phases may form during sintering of SiC/Al4SiC4 with the yttrium (Y) additive. This would significantly promote the densification of SiC/Al4SiC4.
Therefore, PCS and Al powders were used in the present study to form Al4SiC4 ceramics with the in-situ formed SiC particles. 5 wt% Y was used as the sintering additive to enable densification of the SiC/Al4SiC4 composites at relatively low temperatures. The aim of this study was to investigate the densification mechanism of SiC/Al4SiC4 composites, and to understand the effect of the nano-sized SiC on the mechanical properties and thermal conductivity of SiC/Al4SiC4 composites. This is the first study to develop highly dense SiC/Al4SiC4 composites using a rare earth element as the sintering additive. The newly developed SiC/Al4SiC4 composites may combine the individual advantages of SiC and Al4SiC4 ceramics, making the SiC/Al4SiC4 materials of great interest for aerospace and nuclear applications.
2. 1 Materials
The PCS powder with an average molecular weight of ~1500 was used as the raw material. It was synthesized by heating polydimethylsilane (PDMS) at the temperature of 455 ℃ under an external pressure of 16 MPa, according to the approach proposed by Hasegawa et al. [28,29]. The Al powder (99% purity, Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) with a particle size of 75–150 μm was also used as the raw material. In addition, the Y powder (Sinopharm Chemical Reagent Co., Ltd., Shanghai, China) with a purity of 99.5% and a mean particle size of 75 μm was used as the sintering additive. The commercially available graphite powder (300 mesh, 99.5% purity) was obtained from Aladdin Chemical Reagent, China. The β-SiC powder (99.5%, Eno Material Co., Ltd., Qinhuangdao, China) with a mean particle size of 0.5 μm was used to synthesize the reference, pure Al4SiC4 sample.
2. 2 Composite preparation
As-synthesized PCS powder was cross-linked at 600 ℃ for 2 h in a flowing Ar atmosphere. The weight fractions of Si, C, and O in the as-obtained cross-linked PCS were 55.45%, 40.1%, and 4.5%, respectively. The amount of free carbon was calculated to be ~18 wt%, assuming that all oxygen was in the form of SiO2 in the cross-linked PCS powder. The cross-linked PCS powder was subsequently mixed with the Al powder. In order to form Al4C3, the molar ratio of Al:C in the cross-linked PCS powder was 4:3. The powder mixture was then pyrolysized at 1000 ℃ for 30 min in an Ar atmosphere. The sample was referred as S1000. The weight fraction of SiC in the final SiC/Al4SiC4 materials was calculated to be 45.4 wt% (43.9 vol%), assuming that all Al transformed to Al4SiC4.
The amount of 5 wt% Y was added to the S1000 sample as the sintering additive. Sintering of the powder mixture was performed in a spark plasma sintering (SPS) furnace (HPD 25/1, FCT systems, Germany) using an Ar atmosphere. The heating rate was 50 ℃/min up to the temperature of 1400 ℃, while it was decreased to 25 ℃ /min for the temperatures above 1400 ℃. The applied uniaxial pressure was increased from 16 to 64 MPa within the first 5 min at the target temperature, followed by the dwell time of additional 5 min at the target temperature. The materials were sintered at different target temperatures, i.e., 1200 ℃ (S1200), 1400 ℃ (S1400), 1500 ℃ (S1500), 1550 ℃ (S1550), 1600 ℃ (S1600), and 1650 ℃ (S1650). For the sake of comparison, the SiC/Al4SiC4 sample without Y additive was sintered using the same conditions as used for the S1650 sample, and was referred as S1650-C. At the same time, the reference monolithic Al4SiC4 sample was also prepared by the SPS sintering of Al4C3 and SiC powders at the temperature of 1900 ℃ for 20 min with a uniaxial pressure of 32 MPa. The Al4C3 powder was synthesized using pure Al and C powders (with the molar ratio of Al:C = 4:3) at 1200 ℃ for 2 h in an Ar atmosphere.
2. 3 Material characterization
The phase composition of the composites was determined by an X-ray diffractometer (XRD, D8 Advance; Bruker AXS, Karlsrude, Germany) using a Cu Kα radiation (λ = 1.5406 Å). The microstructure and the fracture surfaces of the specimens were performed using a scanning electron microscope (SEM, FEI Quanta FEG 250), equipped with an energy dispersive spectroscopy (EDS) detector.
The apparent densities of the composites were measured using an Archimedes’ method. The porosity of the samples was determined using a High-Performance Full-Automatic Mercury Intrusion Meter (AutoPore IV 9510, micrometrics instrument corporation, USA). The hardness measurements were carried out using a Vickers’ diamond indenter (HVs-1000 Digital micro Vickers’ Hardness Tester, Beijing Times Mountain Peak Technology Co., Ltd., Beijing, China). An indentation load (P) of 9.8 N and a dwell time of 15 s were used. At least 20 indents were measured for each specimen. The elastic modulus was measured using a nanoindentation system (Hysitron PI85, Bruker) on the polished samples. The indentation fracture toughness (KIC) was determined by the measuring of the length of the cracks introduced by Vickers indentation at the load of 9.8 N. The fracture toughness was then calculated using Eq. (1), according to Refs. [30,31]:
KIC = 0.016(E/H)1/2PC–3/2 (1)
where E is the elastic modulus, H is the indentation hardness, and C is the half-length of the cracks, measured using SEM.
The laser flash method was used to measure the thermal diffusivity and specific heat capacity using a Netzsch LFA 457 apparatus (LFA, NETZSCH-Gerätebau GmbH, Germany). The thermal conductivity (κ, W/(m·K)) was calculated according to Eq. (2) :
κ = αρCp (2)
where α is the coefficient of thermal diffusivity, ρ is the density of the material, and Cp is the specific heat capacity.
3 Results and discussion
3. 1 Microstructure of SiC/Al4SiC4 composites and Al4SiC4
The relative densities of the SiC/Al4SiC4 composites sintered at different temperatures are shown in Table 1. The relative density significantly increased from 67.4% to 98.6% when the sintering temperature was increased from 1400 to 1650 ℃ for the sample with 5 wt% Y additive. On the other hand, the relative density was only 89.9% for the sample sintered at 1650 ℃ without Y additive (S1650-C). This clearly indicates that the Y additive promoted densification of the SiC/Al4SiC4 composites, probably due to formation of a liquid phase.
Table 1 Chemical composition, density, grain size of SiC particles, mechanical and thermal properties of the investigated SiC/Al4SiC4 multiphase ceramics and Al4SiC4 monoliths
Figure 1 shows the XRD patterns of the pyrolyzed powder (S1000) and all the samples sintered at different temperatures. The following phases: Al4C3 (JCPDS card No. 35-0799), Al–Si alloy (JCPDS card No. 41-1222), and β-SiC (JCPDS card No. 29-1129) were detected for the pyrolyzed powder (S1000). The broad diffraction peak of the β-SiC phase suggests an incomplete crystallization of SiC at 1000 ℃, according to Reaction (3) . The presence of the Al4C3 phase can be explained by the reaction of molten Al (~660 ℃) with free carbon, according to Reaction (4) . On the other hand, some remaining Al could also react with SiC to form Al4C3 and Si, according to Reaction (5) . The Al–Si alloy could be subsequently formed by the dissolution of Si into the unreacted Al, according to Reaction (6) . The main possible reactions could be summarized as follows:
PCS → SiC + C + gas↑ (3)
4Al + 3C → Al4C3 (4)
4Al + 3SiC → Al4C3 + 3Si (5)
xAl + ySi → AlxSiy (6)
Fig. 1 XRD patterns of the pyrolyzed powder (S1000) and the samples sintered at different temperatures.
When the calcined powder with the Y additive was sintered at 1200 ℃, Al4SiC4 (JCPDS card No. 35-1072), Y3Si2C2 (JCPDS card No. 51-0835), and YAlO3 impurity (JCPDS card No. 33-0041) were detected by XRD. Some residual unreacted SiC (JCPDS card No. 29-1126) and Al4C3 were also found (Fig. 1(a)). It should be mentioned that the β- to α- phase transformation of SiC was observed. This might have been caused by the dissolution of Al into the SiC grains,which promotes the β-to-α phase transformation of SiC [34,35]. The Al4SiC4 phase could be formed at the temperatures above 1106 ℃, according to the pseudo binary phase diagram of the Al4C3–SiC system . The formation of Y3Si2C2 occurred by the reaction between Y and SiC, according to Reaction (8) . There were no new, additional phases detected when the sintering temperature was increased to 1400 ℃. The main possible reactions in the temperature range of 1200–1400 ℃ could be summarized as follows:
Al4C3 + SiC → Al4SiC4 (7)
3Y + 2SiC → Y3Si2C2 (8)
When the sintering temperature increased to 1500–1600 ℃, the diffraction peaks of Y3Si2C2 were no longer observed in the XRD patterns. This was caused by the formation of a liquid phase, which forms at the temperatures around 1560 ℃ via a ternary eutectic Reaction (9), according to the calculated Y–Si–C ternary phase diagram . In Reaction (9), “L” represents the liquid phase and “γ” represents the binary Y–C compound . In addition, the decomposition of Y3Si2C2 led to the formation of Y-based liquid, according to Reaction (10) . The Y-based liquid could later oxidize by its reaction with oxygen from the composites, according to Reaction (11) . Thus, the main possible reactions could be described as follows:
L ⇋ SiC + γ + Y3Si2C2 (9)
Y3Si2C2 → 2SiC + 3Y(l) (10)
4Y(l) + O2 → 2Y2O3 (11)
When the temperature further increased to 1650 ℃, the relative intensities of the SiC diffraction peaks increased significantly. This was caused by the growth of the SiC grains. According to the calculated phase diagram, an eutectic liquid phase is formed in the Y2O3–Al2O3–SiO2 system at the temperatures between 1635 and 1644 ℃ by Reactions (12) and (13) . Furthermore, Liang et al.  reported that a liquid phase was clearly observed in the pressureless sintered SiC with Al2O3 and Y2O3 additives at the temperature of 1680 ℃. The presence of a liquid phase improves the atoms diffusion and promotes the grain growth of SiC. On the other hand, a certain amount of an amorphous silicate phase (could not be detected by XRD) might have been formed by the reactions between the additives and the silica layer present on the surface of SiC, as observed by Falk .
L + Y4Al2O9 ⇋ Y2SiO5 + Y3Al5O12 (12)
L + YAlO3 + Y4Al2O9 ⇋ Y3Al5O12 (13)
Moreover, the relative X-ray intensities of the (001) planes of Al4SiC4 increased with the increasing sintering temperature. The relative ratio of the (00 10) and (11–0) planes (I(00 10)/I(11–0)) of Al4SiC4 increased from 0.9 to 1.5 when the sintering temperature increased from 1500 to 1650 ℃. This relative ratio is much higher than the ratio calculated from the standard JCPDS card (I(00 10)/I(11–0) = 0.76). This indicates that preferential grain growth orientation was obtained. Xing et al.  observed similar phenomena during the synthesis of Al4SiC4 powder by the carbothermal reduction method.
The sample sintered without Y additive (S1650-C) contained the following crystalline phases: SiC, Al4SiC4, and the impurity of Al4O4C. The formation of Al4O4C can be possible explained by the reduction of Al2O3 by its reaction with C (Reaction (14)), as well as to the reaction between Al2O3 and Al4C3 (Reaction (15)) [9, 19].
2Al2O3 + 3C → Al4O4C + 2CO(g) (14)
4Al2O3 + Al4C3 → 3Al4O4C (15)
For the sake of comparison, the monolithic Al4SiC4 was also prepared. Figure 2(a) shows the XRD patterns of both the Al4C3 powder as the raw material, and the as-obtained pure Al4SiC4 ceramics. No impurity was detected. The SEM analysis of the fracture surface revealed the presence only a few micro-pores (Fig. 2(b)), suggesting that a high dense Al4SiC4 material (98.9%) was obtained.
Fig. 2 (a) XRD patterns of the Al4C3 powder and pure Al4SiC4 sintered at 1900 ℃; (b) SEM image of the fracture surface of pure Al4SiC4 ceramics.
To further understand the effect of Y additive on the sintering of SiC/Al4SiC4 composites and reveal the densification mechanism, the fracture surfaces of the samples sintered at different temperatures (Figs. 3 and 4) were observed. Obviously, the microstructure of the S1400 sample was significantly porous, and the agglomeration of SiC particles was also observed (Figs. 3(a) and 4(a)). The number of pores significantly decreased when the sintering temperature increased to 1500 ℃ (Fig. 3(b)). No obvious change in number of pores was observed when the sintering temperature was increased to 1550 and 1600 ℃ (Figs. 3(c) and 3(d)). However, some abnormal growth of Al4SiC4 grains was observed at these temperatures (Figs. 4(c) and 4(d)). The pores were always located around the large Al4SiC4 grains.
Fig. 3 Low magnification SEM images of the fracture surfaces of the SiC/Al4SiC4 sintered at different temperatures: (a) S1400, (b) S1500, (c) S1550, (d) S1600, (e) S1650, and (f) S1650-C (without Y additive).
Fig. 4 High magnification SEM images of the fracture surfaces of the SiC/Al4SiC4 sintered at different temperatures: (a) S1400, (b) S1500, (c) S1550, (d) S1600, (e) S1650, and (f) S1650-C (without Y additive).
A significantly different morphology of the fracture surface was observed for the specimen S1650 (Fig. 3(e)). The SEM analysis revealed the presence of a rough surface, resulting from the grains pull out. No obvious porosity was observed. On the other hand, the fracture surface of the S1650-C sample sintered without Y additive was relatively smooth (Figs. 3(f) and 4(f)). Numerous pores and abnormal growth of Al4SiC4 grains were observed. Such a significantly different fracture surface morphology was caused by a large amount of the liquid phase formed in the temperature range of 1600–1650 ℃ when Y sintering additive was used (the sample S1650) [27,37]. The large amount of the liquid phase not only filled up the pores between the large plate-like Al4SiC4 grains (Figs. 3(e) and 4(e)), but also improved the diffusion of atoms along the grain boundaries [39,40]. Therefore, the relative density of the S1650 (98.6%) was significantly higher than that of the S1650-C (89.9%).
Figure 5 shows the typical low and high magnification back-scattered electron (BSE) SEM images of the samples S1500, S1600, and S1650. In the case of the samples sintered at 1500 and 1600 ℃, the Y-rich amorphous glassy phase (bright phase) formed by solidification of a liquid upon cooling from the sintering temperature, was localized in the porous area between the large Al4SiC4 grains (Figs. 5(a), 5(b), 5(d), and 5(e)). However, the liquid phase was homogenously dispersed when the sintering temperature was increased to 1650 ℃. Most of the Y-rich amorphous phase was located at the grain boundaries (Figs. 5(c) and 5(f)). The uniform dispersion of the glassy phase at the grain boundaries would play an important role for the mechanical properties of the composites.
Fig. 5 Typical low and high magnification BSE SEM images of the fracture surfaces of SiC/Al4SiC4 showing the distribution of the Y-rich phase sintered at: (a, d) 1500 ℃; (b, e) 1600 ℃; (c, f) 1650 ℃.
3. 2 In-situ reaction and densification process
According to the phase and microstructure evolution, the in-situ reaction and sintering process of SiC/Al4SiC4 composites can be divided into four stages, as schematically shown in Fig. 6. In the initial stage (≤ 1400 ℃), the Al4SiC4 and Y3Si2C2 phases were formed via Reactions (7) and (8). There was no liquid phase formed, so any densification occurred via solid-state mechanism. The formation of the large, interconnected pores in the sintered body was characteristic for this stage. As the sintering temperature increased to 1500 ℃, Y3Si2C2 transformed to a liquid phase by Reaction (9). The liquid phase wetted and covered the SiC particles. The grains of SiC and Al4SiC4 dissolved in the liquid phase, resulting in the particles rearrangement. In the temperature range of 1500–1600 ℃, the sintering occurred via solution and re-precipitation process. However, the densification was limited by the formation of a small amount of the liquid phase at this temperature range.
Fig. 6 Schematic diagram of the in-situ reaction and densification process of SiC/Al4SiC4 composites at different sintering temperatures.
When the sintering temperature increased to 1650 ℃, a large amount of the liquid phase was formed by Reactions (12) and (13). The sufficient amount of the liquid phase not only filled the pores, but also promoted the diffusion of atoms along the grain boundaries . Furthermore, the coarsening of the SiC and Al4SiC4 grains occurred via the solution-reprecipitation mechanism. At high temperatures, the solubility of C, Si, and Al in the liquid phase increases, due to the large capillary forces. Thus, the fine SiC and Al4SiC4 grains dissolved into the liquid, and reprecipitated onto the surface of larger SiC and Al4SiC4 grains. The initial grains of SiC and Al4SiC4 acted as nuclei for the growth of large grains via Oswald ripening mechanism . Therefore, the mean grain size of the SiC particles significantly increased from 153 to 304 nm when the temperature increased from 1600 to 1650 ℃.
3. 3 Mechanical properties
Table 1 shows the Vickers hardness, elastic modulus, and KIC of the SiC/ Al4SiC4 composites sintered at different temperatures. For the sake of comparison, the mechanical properties of the monolithic Al4SiC4 material are also given in Table 1. The Vickers hardness increased from 15.3 ± 1.9 (S1600) to 16.3 ± 0.6 GPa (S1650), as the density increased with the increasing sintering temperature. However, the hardness of the sample without Y additive (S1650-C) was higher than that of the S1650 sample, even though the S1650-C sample exhibited a lower relative density than the S1650 sample. This can be attributed to the finer grain size of the SiC particles in the S1650-C (131 nm) sample, when compared to the S1650 sample (304 nm). Another reason of the lower hardness of the S1650 sample was that it contained a relatively large amount of the grain boundary glassy phase, which has a significantly lower hardness than the crystalline ceramic grains . The hardness of the SiC/Al4SiC4 composite sintered at 1650 ℃ (16.3 GPa) was significantly higher when compared to the pure Al4SiC4 material (13.2 GPa). This can be explained by the presence of SiC particles with a relatively high intrinsic hardness (25 GPa ). Since the open porosity of the S1650 sample (6.5%) was at the similar level than that of the pure Al4SiC4 material (5.6%), the effect of porosity/density on the hardness results can be neglected.
The elastic modulus of the sample S1650 was slightly lower than that of the S1600 and S1650-C samples (Table 1). This was probably caused by the presence of a large amount of the secondary phases with low elastic modulus in the S1650 sample. The amorphous silicate-like phases possess a low elastic modulus, which would decrease the elastic modulus of the whole composite body .
Table 1 shows the fracture toughness of the SiC/Al4SiC4 composites sintered at different temperatures. The fracture toughness of the composite sintered with Y additive at 1650 ℃ (4.9 ± 0.2 MPa·m1/2) was significantly higher than the fracture toughness of both the composite sintered without Y additive (2.7 ± 0.09 MPa·m1/2) and the pure Al4SiC4 monolith (2.2 ± 0.2 MPa·m1/2). Such a relatively high fracture toughness of the S1650 sample can be attributed mainly to its high density, and the presence of the fine SiC particles. Furthermore, the chemistry of the grain boundary phases would also play an important role in toughening, as the weak interface could improve fracture toughness. In order to understand the fracture mechanisms, the propagation of the indentation cracks and the fractured surfaces of the composites sintered at 1650 ℃ (with and without Y additive) were observed by SEM.
Figures 7(a) and 7(c) show the BSE images of the indentation crack paths in the S1650 and S1650-C samples. The straight crack path and mainly transgranular fracture were observed in the S1650-C sample (Fig. 7(c)). The corresponding fracture surface was relatively smooth (Fig. 7(d)). Transgranular failure was the main fracture mechanism, which suggests a low fracture toughness of the material.
Fig. 7 Typical SEM microstructure images of (a, b) S1650 and (c, d) S1650-C samples: (a, c) the crack propagation path after the indentation test; (b, d) the fracture surfaces.
On the other hand, some typical toughening mechanisms, such as crack deflection and crack bridging, were observed in the sample S1650 (Fig. 7(a)). The crack bridging and crack deflection inhibited the crack growth, resulting in a high fracture toughness of the material. The corresponding fracture surface was relatively rough (Fig. 7(b)), consisting of both the intergranular and transgranular failure modes. The intergranular fracture should constitute a major contribution to the high fracture toughness obtained for the material . The intergranular fracture of S1650 was indicated by the holes as shown in Fig. 7(b). The holes were formed by the pullout of the SiC and some Al4SiC4 grains, indicating a relatively low interfacial strength at the grain boundaries . Such a weak interface led to the more tortuous crack path and the intergranular fracture in the SiC/Al4SiC4 composite sintered with Y additive at 1650 ℃.
The SEM analysis also revealed that the size of the Al4SiC4 grains was significantly different than that of the SiC grains, leading to a broad grain size (bimodal) distribution of the composite. In general, the bimodal grain size distribution is beneficial for the increase in the number of crack deflection events and other toughening mechanisms . The crack deflection consumed more fracture energy, resulting in a higher fracture toughness of the material.
3. 4 Thermal properties
Figure 8 shows the thermal conductivity of the SiC/Al4SiC4 composites sintered at different temperatures. The thermal conductivity increased with the increasing sintering temperature. The thermal conductivity of the S1500 sample was only 5.6 W/(m·K), whereas it significantly increased to 15.8 W/(m·K) for the specimen S1650. On the other hand, the thermal conductivities of the monolithic Al4SiC4 material and the composite without Yadditive (S1650-C) were 7.8 and 9.6 W/(m·K), respectively. Thus, the thermal conductivity of Al4SiC4 was significantly improved by the incorporation of the fine grained SiC with high intrinsic thermal conductivity, as well as by the beneficial effect of the Y sintering additive.
Fig. 8 Thermal conductivity of the SiC/Al4SiC4 composites sintered at different temperatures with and without Y sintering additive (S1650-C).
The thermal conductivity of ceramics is determined by the density, grain size, as well as phase composition of the specimens. The S1650 specimen showed the highest density among all of the investigated materials. Moreover, the heat in the SiC/Al4SiC4 composites is mainly transferred by phonons, due to the lack of electronic conductivity of SiC and Al4SiC4 . The material with a smaller grain size contains more grain boundaries per volume unit, which results in the phonon scattering at the grain boundaries. This significantly decreases the thermal conductivity of the material. The S1650 sample contained much larger SiC grains (304 ± 33 nm) when compared to the samples of S1600 (153 ± 25 nm) and S1650-C (131 ± 16 nm), which can well explain the higher thermal conductivity of the S1650 sample. It should also be mentioned that there was obviously more liquid phase formed at the grain boundaries and grains junctions in the S1650 sample, when compared to the samples S1600 and S1650-C. A higher amount of the grain boundary phases may be harmful to the thermal conductivity of the materials. In the present study, however, the large grain size of SiC and the high density of the S1650 sample obviously eliminated the negative effect of the higher fraction of the grain boundary phases. It can be concluded that the highest thermal conductivity of the SiC/Al4SiC4 composite sintered with Y additive at 1650 ℃ was attributed to its high relative density and the relatively large grain size, resulting in less phonon scattering.
The highly dense SiC/Al4SiC4 composites (98.6%) with the improved hardness (16.3 ± 0.6 GPa), fracture toughness (4.9 ± 0.2 MPa·m1/2), and thermal conductivity (15.8 W·m–1·K–1) were fabricated by SPS at a relatively low temperature of 1650 ℃. The materials were obtained by the in-situ reaction sintering of the starting mixture of PCS, Al, and 5 wt% Y sintering additive. Both, the phase transformation of Y3Si2C2 at the temperatures around 1560 ℃, and the liquid phase formation in the SiC(SiO2)–Al2O3–Y2O3 systems at around 1650 ℃ significantly promoted the densification of SiC/Al4SiC4 composites. The in-situ reactions led to the formation of fine grained SiC particles in the microstructure. The hardness, elastic modulus, fracture toughness, and thermal conductivity of Al4SiC4 ceramics were significantly improved by the incorporation of the SiC particles. The improved mechanical properties and thermal conductivity of the SiC/Al4SiC4 composite sintered at 1650 ℃ were attributed to its high density, appropriate grain size and bimodal grain size distribution, as well as to the appropriate grain boundary structure.
 Liao T, Wang JY, Zhou YC. Atomistic deformation modes and intrinsic brittleness of Al4SiC4: A first-principles investigation. Phys Rev B 2006, 74: 174112.
 Xing XM, Li B, Chen JH, et al. Formation mechanism of large size plate-like Al4SiC4 grains by a carbothermal reduction method. CrystEngComm 2018, 20: 1399–1404.
 Sun L, Gao YM, Yoshida K, et al. Prediction on structural, mechanical and thermal properties of Al4SiC4, Al4C3 and 4H-SiC under high pressure by first-principles calculation. Mod Phys Lett B 2017, 31: 1750080.
 Xing XM, Chen JH, Bei GP, et al. Synthesis of Al4SiC4 powders via carbothermic reduction: Reaction and grain growth mechanisms. J Adv Ceram 2017, 6: 351–359.
 Le-Tran HL, Sarigiannidou E, Gélard I, et al. Vaporization and condensation in the Al4C3-SiC system. J Eur Ceram Soc 2017, 37: 4475–4482.
 Chen JH, Zhang ZH, Mi WJ, et al. Fabrication and oxidation behavior of Al4SiC4 powders. J Am Ceram Soc 2017, 100: 3145–3154.
 Huang XX, Wen GW, Cheng XM, et al. Oxidation behavior of Al4SiC4 ceramic up to 1700 ℃. Corros Sci 2007, 49: 2059–2070.
 Yamamoto O, Ohtani M, Sasamoto T. Preparation and oxidation of Al4SiC4. J Mater Res 2002, 17: 774–778.
 Wen GW, Huang XX. Increased high temperature strength and oxidation resistance of Al4SiC4 ceramics. J Eur Ceram Soc 2006, 26: 1281–1286.
 Yu C, Deng CJ, Zhu HX, et al. Slag corrosion resistance of Al4SiC4. Int J Mater Res 2017, 108: 249–252.
 Yu C, Cheng KR, Ding J, et al. Synthesis and some properties of Al4SiC4-Al4O4C composites. Ceram Int 2018, 44: 17154–17159.
 Magnani G, Galvagno S, Sico G, et al. Sintering and mechanical properties of β-SiC powder obtained from waste tires. J Adv Ceram 2016, 5: 40–46.
 Fu CL, Yang Y, Huang ZR, et al. Investigation on the laser ablation of SiC ceramics using micro-Raman mapping technique. J Adv Ceram 2016, 5: 253–261.
 Wang HL, Gao ST, Peng SM, et al. KD-S SiCf/SiC composites with BN interface fabricated by polymer infiltration and pyrolysis process. J Adv Ceram 2018, 7: 169–177.
 Liu GW, Zhang XZ, Yang J, et al. Recent advances in joining of SiC-based materials (monolithic SiC and SiCf/SiC composites): Joining processes, joint strength, and interfacial behavior. J Adv Ceram 2019, 8: 19–38.
 Zhou YC, Wan DT, Bao YW, et al. In situ processing and high-temperature properties of Ti3Si(Al)C2/SiC composites. Int J Appl Ceram Technol 2006, 3: 47–54.
 Wan DT, Zhou YC, Bao YW, et al. In situ reaction synthesis and characterization of Ti3Si(Al)C2/SiC composites. Ceram Int 2006, 32: 883–890.
 Zhang JF, Wu T, Wang LJ, et al. Microstructure and properties of Ti3SiC2/SiC nanocomposites fabricated by spark plasma sintering. Compos Sci Technol 2008, 68: 499–505.
 Inoue K, Yamaguchi A. Synthesis of Al4SiC4. J Am Ceram Soc 2003, 86: 1028–1030.
 Lee SH, Oh HC, An BH, et al. Ultra-low temperature synthesis of Al4SiC4 powder using spark plasma sintering. Scr Mater 2013, 69: 135–138.
 Hasegawa M, Itatani K, Aizawa M, et al. Low-temperature synthesis of aluminum silicon carbide using ultrafine aluminum carbide and silicon carbide powders. J Am Ceram Soc 1996, 79: 275–278.
 Oscroft RJ, Thompson DP. Influence of oxygen on the formation of aluminum silicon carbide. J Am Ceram Soc 1992, 75: 224–226.
 Lee JS, Lee SH, Nishimura T, et al. Hexagonal plate-like ternary carbide particulates synthesized by a carbothermal reduction process: Processing parameters and synthesis mechanism. J Am Ceram Soc 2009, 92: 1030–1035.
 Yu C, Yuan WJ, Deng CJ, et al. Synthesis of hexagonal plate-like Al4SiC4 from calcined bauxite, silica and carbon black. Powder Technol 2013, 247: 76–80.
 Shao JQ, Li M, Chang KK, et al. Fabrication and characterization of SPS sintered SiC-based ceramic from Y3Si2C2-coated SiC powders. J Eur Ceram Soc 2018, 38: 4833–4841.
 Zhou YC, Xiang HM, Dai FZ. Y5Si3C and Y3Si2C2: Theoretically predicted MAX phase like damage tolerant ceramics and promising interphase materials for SiCf/SiC composites. J Mater Sci Technol 2019, 35: 313–322.
 Pan Z, Fabrichnaya O, Seifert HJ, et al. Thermodynamic evaluation of the Si-C-Al-Y-O system for LPS-SiC application. J Phase Equilib Diffus 2010, 31: 238–249.
 Hasegawa Y, Okamura K. Synthesis of continuous silicon carbide fibre. J Mater Sci 1983, 18: 3633–3648.
 Birot M, Pillot JP, Dunogues J. Comprehensive chemistry of polycarbosilanes, polysilazanes, and polycarbosilazanes as precursors of ceramics. Chem Rev 1995, 95: 1443–1477.
 Anstis GR, Chantikul P, Lawn BR, et al. A critical evaluation of indentation techniques for measuring fracture toughness: I, direct crack measurements. J Am Ceram Soc 1981, 64: 533–538.
 Unlu MD, Goller G, Yucel O, et al. Processing and mechanical characterisation of monolithic silicon carbide ceramic consolidated by spark plasma sintering (SPS). Int J Mater Res 2013, 104: 1240–1246.
 Volz E, Roosen A, Hartung W, et al. Electrical and thermal conductivity of liquid phase sintered SiC. J Eur Ceram Soc 2001, 21: 2089–2093.
 Carotenuto G, Gallo A, Nicolais L. Degradation of SiC particles in aluminium-based composites. J Mater Sci 1994, 29: 4967–4974.
 Yu C, Wu ZM, Ding J, et al. Effect of Al4SiC4 additive on the fabrication and characterization of recrystallized SiC honeycomb ceramics. Ceram Int 2019, 45: 16612–16617.
 Lee JS, Ahn YS, Nishimura T, et al. Ultra-low-temperature sintering of nanostructured β-SiC. J Am Ceram Soc 2011, 94: 324–327.
 Xu K, Zou HK, Chang KK, et al. Thermodynamic description of the sintering aid system in silicon carbide ceramics with the addition of yttrium. J Eur Ceram Soc 2019, 39: 4510–4519.
 Liang HQ, Yao XM, Huang ZR, et al. In situ analysis of the liquid phase sintering process of alpha-SiC ceramics. J Inorg Mater 2016, 31: 443.
 Falk LKL. Microstructural development during liquid phase sintering of silicon carbide ceramics. J Eur Ceram Soc 1997, 17: 983–994.
 German RM, Suri P, Park SJ. Review: Liquid phase sintering. J Mater Sci 2009, 44: 1–39.
 Sciti D, Guicciardi S, Bellosi A. Effect of annealing treatments on microstructure and mechanical properties of liquid-phase-sintered silicon carbide. J Eur Ceram Soc 2001, 21: 621–632.
 El-Atwani O, Quach DV, Efe M, et al. Multimodal grain size distribution and high hardness in fine grained tungsten fabricated by spark plasma sintering. Mater Sci Eng: A 2011, 528: 5670–5677.