Silicon carbide ceramic with high thermal conductivity and bending strength and preparation method thereof

Abstract

The application provides a silicon carbide ceramic with high thermal conductivity and bending strength and a preparation method thereof. The silicon carbide ceramic is prepared by using a Re2O3-ReF3-Bi2O3 (Re=La, Y, Yb, Sc) ternary system as a sintering aid; the preparation method of the silicon carbide ceramic has the advantages of low sintering temperature, high density, small grain size, high thermal conductivity, excellent mechanical properties and adjustable thermal-mechanical performance; when the mass fraction of SiC powder in the mixed powder is 94.7%, the sum of the mass fractions of Re2O3 and ReF3 in the mixed powder is 5%, the mass fraction of Bi2O3 is 0.3%, and Re is Y, the liquid phase content of the 5YF-BiO sample prepared is appropriate, the densification is complete, the interface structure is fine, the thermal performance (thermal conductivity 197 W / (m*K)) is optimized, the bending strength (416 MPa) is considered at the same time, and the comprehensive performance is better.

Description

Technical Field

[0001] This invention relates to the field of silicon carbide ceramic preparation technology, and in particular to a silicon carbide ceramic with both high thermal conductivity and flexural strength, and a method for preparing the same. Background Technology

[0002] Silicon carbide (SiC) ceramics possess excellent properties such as high hardness, high strength, high temperature resistance, corrosion resistance, low coefficient of thermal expansion, and high thermal conductivity, making them promising for applications in semiconductor equipment, high-temperature structural components, heat exchangers, nuclear industry, and aerospace. However, silicon carbide is a strongly covalent compound with an extremely low atomic self-diffusion coefficient, making it difficult to densify pure silicon carbide powder through conventional sintering. Therefore, sintering aids are typically introduced, and high-temperature sintering processes are used for its preparation.

[0003] Currently, the main sintering methods for silicon carbide ceramics include pressureless sintering, hot pressing, and reaction sintering. Among these, liquid-phase sintering of silicon carbide ceramics has become a research hotspot due to its relatively low sintering temperature, high density, and strong designability. Liquid-phase sintering typically uses oxides as sintering aids, such as binary systems like Y₂O₃-Al₂O₃, to promote densification.

[0004] However, existing methods for preparing silicon carbide ceramics still have the following drawbacks: First, achieving both high density and high thermal conductivity is difficult. To increase density, it's often necessary to increase the amount of sintering aids or raise the sintering temperature. Increasing the amount of sintering aids leads to an increase in the content of grain boundary glass phase, and the thermal conductivity of the grain boundary phase is much lower than that of silicon carbide grains, thus significantly reducing the overall thermal conductivity of the ceramic. High temperatures not only lead to a significant increase in energy consumption but also cause abnormal growth of silicon carbide grains, reducing the mechanical strength and fracture toughness of the ceramic. How to achieve high thermal conductivity while maintaining high density is a contradiction that current technologies struggle to resolve.

[0005] Secondly, the control of thermo-mechanical properties is challenging. As the application of silicon carbide ceramics in high-end equipment continues to expand, higher demands are placed on the comprehensive properties of materials, such as thermal conductivity, flexural strength, fracture toughness, and hardness. Existing binary additive systems struggle to achieve an optimal balance between high thermal conductivity and high strength simultaneously, often resulting in a trade-off—pursuing high thermal conductivity at the expense of strength, or pursuing high strength at the expense of insufficient thermal conductivity. How to achieve synergistic optimization and precise control of thermo-mechanical properties through the design of sintering additive systems is a pressing technical challenge that needs to be addressed in this field.

[0006] Third, the poor stability of grain boundary phases limits high-temperature performance. The grain boundary glass phases formed by existing oxide additives are prone to softening, crystallization, or volatilization at high temperatures, leading to a sharp decline in the mechanical properties and thermal conductivity of the ceramics under high-temperature conditions, thus limiting the application of silicon carbide ceramics in high-temperature environments. Especially for high-temperature hot components in semiconductor manufacturing equipment, the high-temperature stability of the grain boundary phase directly affects the lifespan and reliability of the devices.

[0007] Therefore, developing a method for preparing silicon carbide ceramics with low sintering temperature, high density, high thermal conductivity, excellent mechanical properties, and tunable thermo-mechanical properties is of great technical significance and application value. Summary of the Invention

[0008] To address the problems existing in the prior art, this invention provides a silicon carbide ceramic with both high thermal conductivity and flexural strength, and a method for preparing the same. This invention employs spark plasma sintering technology, using SiC powder as raw material and a Re2O3-ReF3-Bi2O3 (Re = La, Y, Yb, Sc) ternary system as a sintering aid to prepare a densified silicon carbide ceramic. This solves the problems of high sintering temperature, difficulty in achieving both thermal conductivity and mechanical properties, and poor grain boundary phase stability in existing silicon carbide ceramics.

[0009] To achieve the above objectives, the present invention adopts the following technical solution: In a first aspect, the present invention provides a method for preparing silicon carbide ceramics that possess both high thermal conductivity and flexural strength, comprising the following steps: SiC powder and ternary sintering aid are mixed to obtain a mixed powder. The mixed powder was subjected to spark plasma sintering to obtain silicon carbide ceramics; The ternary sintering aids include Re2O3, ReF3 and Bi2O3, wherein Re is selected from any one of La, Y, Yb and Sc.

[0010] Preferably, the mass fraction of SiC powder in the mixed powder is 92.7-96.7%, the sum of the mass fractions of Re2O3 and ReF3 in the mixed powder is 3-7%, and the remainder is Bi2O3.

[0011] Preferably, the molar ratio of Re2O3 to ReF3 is (1~2):(1~2).

[0012] Preferably, the mixed powder is placed in a spark plasma sintering furnace, subjected to a pressure of 38-40 MPa under a vacuum atmosphere, and heated to 1850-1950°C, and held at this temperature for 10-15 minutes to obtain silicon carbide ceramic.

[0013] Preferably, SiC powder and ternary sintering aid are mixed and then ethanol is added, followed by ball milling and drying to obtain mixed powder.

[0014] Preferably, SiC powder and ternary sintering aid are mixed and then added to ethanol and placed in a ball mill jar. Grinding balls are added, and the mixture is ball-milled at 150~200 r / min for 10~15 h. After drying, the mixed powder is obtained.

[0015] Preferably, the ratio of the mass of the grinding ball to the sum of the masses of SiC powder and ternary sintering aid is (4~6):1.

[0016] Preferably, the ratio of the mass of ethanol to the sum of the masses of SiC powder and ternary sintering aid is (1~2):(1~2).

[0017] Preferably, SiC powder and ternary sintering aid are mixed and then added to ethanol and placed in a ball mill jar. Grinding balls are added, and the mixture is ball-milled at 150~200 r / min for 10~15 h. After drying at 80~100℃ for 20~24 h, the mixture is pulverized and passed through an 80~100 mesh sieve to obtain the mixed powder.

[0018] Secondly, the present invention also provides a silicon carbide ceramic with both high thermal conductivity and flexural strength, which is prepared by the aforementioned preparation method.

[0019] The silicon carbide ceramic of the present invention, which possesses both high thermal conductivity and flexural strength, and its preparation method, have the following advantages over the prior art: 1. This invention relates to a method for preparing silicon carbide ceramics with both high thermal conductivity and flexural strength. The method employs spark plasma sintering (SPCS) technology, using SiC powder as raw material and a Re₂O₃-ReF₃-Bi₂O₃ (Re = La, Y, Yb, Sc) ternary system as a sintering aid to prepare densified silicon carbide ceramics. The type of rare earth element significantly affects the densification behavior, phase composition, and microstructure of SiC ceramics. The Yb and Sc system achieves high densification due to its small ionic radius, high liquid phase viscosity, and good high-temperature stability; while the La and Y system exhibits lower densification due to its low liquid phase viscosity and significant volatilization. The fluorine-oxygen exchange reaction suppresses the adverse effects of aid volatilization. The Y and Yb systems simultaneously generate rare earth silicate phases. Based on this, further research using Y-based aids reveals that the content of Y-based aids exhibits a typical optimal ratio for the densification behavior of SiC ceramics. At 3 wt.%, the liquid phase is insufficient, resulting in incomplete densification. With 5 wt.% of additives, the liquid phase content was moderate, achieving complete densification with a density of 99.1% and an open porosity of 0.12%. However, when the additive content increased to 7 wt.%, the liquid phase was excessive, and volatilization led to a decrease in density. 2. In the silicon carbide ceramics of this invention, the rare earth element types significantly affect the thermal properties of SiC ceramics by regulating liquid phase characteristics and grain growth behavior. The Y system, due to its largest grain size and fewest grain boundaries, achieves the highest thermal conductivity, 170.7 W / (m·K); grain size and grain boundary characteristics are the core factors determining thermal conductivity. Studies on the content of Y-based additives revealed that grain size and porosity significantly affect thermal conductivity. The 5YF-BiO sample, due to its coarse grains and extremely low porosity, achieves a thermal conductivity of 197.1 W / (m·K). Although the 3YF-BiO sample has relatively large grains, the scattering effect of numerous pores limits its thermal conductivity to 170.7 W / (m·K). The 7 wt.% sample, due to grain refinement and increased porosity, has a thermal conductivity reduced to 160.2 W / (m·K). 3. The strength difference of the silicon carbide ceramics of the present invention is the result of the synergistic effect of porosity, grain size, and grain boundary bonding: the Yb system achieves fine grain strengthening through high density, fine grain structure, and uniformly distributed grain boundary phases; although the strength of the Y system is slightly lower than that of the Yb system, it still relies on the strong bonding of grain boundaries to offset the adverse effects of coarse grains and porosity; although the Sc system is extremely dense and has the finest grains, the presence of residual second phase weakens the grain boundary bonding, limiting further improvement in strength; the La system suffers from strength degradation due to stress concentration dominated by a large number of pores. The effect of Y-based additive content on performance shows that: although the 3YF-BiO sample has high porosity, the grain boundaries are clean and the bonding is strong; although the 7YF-BiO sample has the smallest grain size, the strengthening effect is weakened due to the presence of a large number of pores, resulting in the lowest flexural strength; the 5YF-BiO sample has an appropriate liquid phase content, complete densification, and a fine interface structure, which optimizes thermal properties while taking into account flexural strength (416 MPa); overall, the 5YF-BiO sample has better comprehensive performance. Attached Figure Description

[0020] To more clearly illustrate the technical solutions in the embodiments of the present invention or the prior art, the drawings used in the description of the embodiments or the prior art will be briefly introduced below. Obviously, the drawings described below are only some embodiments of the present invention. For those skilled in the art, other drawings can be obtained based on these drawings without creative effort.

[0021] Figure 1 The XRD diffraction patterns of SiC ceramics prepared by different rare earth element additive systems in Example 1 are shown. Figure 2 Scanning electron microscope images and grain size statistics of SiC ceramics prepared by different rare earth element additive systems in Example 1; Figure 3The images shown are STEM, HAADF, HRTEM, and Fast Fourier Transform (FFT) images of the SiC ceramic 3Yb2O3-YF3-Bi2O3 sample prepared in Example 1. Figure 4 SEM images and grain size distribution of silicon carbide ceramic microstructures prepared using Y2O3-YF3-Bi2O3 as a sintering aid in Example 2. Detailed Implementation

[0022] The present invention will now be clearly and completely described with reference to specific embodiments. Obviously, the described embodiments are only a part of the embodiments of the present invention, and not all of them. Based on the embodiments of the present invention, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of the present invention.

[0023] To better understand the invention and not to limit its scope, all figures indicating amounts, percentages, and other numerical values ​​used in this application should, in all cases, be understood to be modified by the word "approximately." Therefore, unless specifically stated otherwise, the numerical parameters listed in the specification and appended claims are approximate values ​​and may vary depending on the desired properties being sought. Each numerical parameter should at least be considered as obtained based on reported significant figures and through conventional rounding methods.

[0024] It should be noted that the order of description of the following embodiments is not intended to limit the preferred order of embodiments. Furthermore, in the description of this application, the term "comprising" means "including but not limited to". Various embodiments of the present invention may exist in the form of a range; it should be understood that the description in the form of a range is merely for convenience and brevity and should not be construed as a rigid limitation on the scope of the invention; therefore, it should be considered that the range description has specifically disclosed all possible sub-ranges and single numerical values ​​within that range. For example, it should be considered that the range description from 1 to 6 has specifically disclosed sub-ranges, such as from 1 to 3, from 1 to 4, from 1 to 5, from 2 to 4, from 2 to 6, from 3 to 6, etc., and single digits within the range, such as 1, 2, 3, 4, 5, and 6, regardless of the range. Additionally, whenever a numerical range is referred to herein, it means including any referenced number (fraction or integer) within the referred range.

[0025] This invention provides a method for preparing silicon carbide ceramics with both high thermal conductivity and flexural strength, comprising the following steps: S1. After mixing SiC powder and ternary sintering aid, a mixed powder is obtained; S2. The mixed powder is subjected to spark plasma sintering to obtain silicon carbide ceramics; The ternary sintering aids include Re2O3, ReF3 and Bi2O3, wherein Re is selected from any one of La, Y, Yb and Sc.

[0026] The method for preparing silicon carbide ceramics of the present invention employs spark plasma sintering technology, using SiC powder as raw material and a Re2O3-ReF3-Bi2O3 (Re = La, Y, Yb, Sc) ternary system as sintering aid to prepare densified silicon carbide ceramics.

[0027] In some embodiments, the mass fraction of SiC powder in the mixed powder is 92.7-96.7%, the sum of the mass fractions of Re2O3 and ReF3 in the mixed powder is 3-7%, and the remainder is Bi2O3. Preferably, the mass fraction of Bi2O3 in the mixed powder is 0.3%. Further preferably, the mass fraction of SiC powder in the mixed powder is 94.7%, the sum of the mass fractions of Re2O3 and ReF3 in the mixed powder is 5%, the mass fraction of Bi2O3 is 0.3%, and Re is Y.

[0028] In some embodiments, the molar ratio of Re2O3 to ReF3 is (1~2):(1~2), and preferably, the molar ratio of Re2O3 to ReF3 is 1:1.

[0029] In some embodiments, the mixed powder is placed in a spark plasma sintering furnace, subjected to a pressure of 38-40 MPa under a vacuum atmosphere, and heated to 1850-1950°C, and held at this temperature for 10-15 min to obtain silicon carbide ceramics. Preferably, a pressure of 40 MPa is applied under a vacuum atmosphere, and the powder is heated to 1900°C, and held at this temperature for 10 min to sinter, thereby obtaining silicon carbide ceramics.

[0030] In some embodiments, SiC powder and ternary sintering aid are mixed and then ethanol is added, followed by ball milling and drying to obtain a mixed powder.

[0031] In some embodiments, SiC powder and ternary sintering aid are mixed and then added to ethanol and placed in a ball mill jar. Grinding balls are added, and the mixture is ball-milled at 150-200 r / min for 10-15 h. After drying, the mixed powder is obtained.

[0032] In some embodiments, the ratio of the mass of the grinding balls to the sum of the masses of SiC powder and ternary sintering aids (i.e., the ball-to-material ratio) is (4~6):1.

[0033] In some embodiments, the ratio of the mass of ethanol to the sum of the masses of SiC powder and ternary sintering aid is (1~2):(1~2).

[0034] In some embodiments, SiC powder and ternary sintering aid are mixed and then added to ethanol and placed in a ball mill jar. Grinding balls are added, and the mixture is ball-milled at 150-200 r / min for 10-15 h. After drying at 80-100℃ for 20-24 h, the mixture is pulverized and passed through an 80-100 mesh sieve to obtain the mixed powder.

[0035] Based on the same inventive concept, the present invention also provides a silicon carbide ceramic, which is prepared by the above-described preparation method.

[0036] The following specific embodiments further illustrate the silicon carbide ceramics and their preparation method of the present invention. This section further explains the content of the present invention in conjunction with specific embodiments, but should not be construed as limiting the present invention. Unless otherwise specified, the technical means used in the embodiments are conventional means well known to those skilled in the art. Unless otherwise specified, the reagents, methods, and equipment used in the present invention are conventional reagents, methods, and equipment in the art.

[0037] In the following examples, the SiC powder was purchased from Shanghai Aladdin Biochemical Technology Co., Ltd., D 50 =0.6 μm; La2O3 powder, Y2O3 powder, Yb2O3 powder, and Sc2O3 powder were purchased from the National Engineering Research Center for Rare Earth Metallurgy and Functional Materials, Ruike; LaF3 powder, YF3 powder, YbF3 powder, and ScF3 powder were purchased from Shanghai Aladdin Biochemical Technology Co., Ltd.; Bi2O3 powder was purchased from Ningxia Juntao New Material Technology Co., Ltd.

[0038] In the following examples, density was determined using the Archimedes method, thermal properties were measured using the laser scintillation method, and flexural strength was tested using the three-point bending test method. The ceramic samples were cut into 10×10×3 mm cubes using a diamond wire cutter. Thermal diffusivity and specific heat capacity were determined using a laser flash analyzer. Prior to measurement, the samples were ultrasonically cleaned and dried, and both surfaces of each sample were coated with a thin layer of graphite to minimize laser reflection and enhance infrared signal sensitivity. Measurements were performed at room temperature. Thermal conductivity was calculated using the following formula.

[0039]

[0040] Where λ is the thermal conductivity of the test sample, W / (m·K); α is the thermal diffusivity of the test sample, mm. 2 / s; ρ is the test value. Sample density, g / cm³ 3 ; C P The specific heat capacity of the test sample is given in J / (g·K).

[0041] The phonon mean free path (L) is defined as the average distance traveled by a phonon during its free motion between two scattering events, and is calculated according to the following formula:

[0042] In the formula, The velocity is phonon velocity. The sound velocity of the SiC ceramic sample used in this calculation is 11820 m / s at room temperature.

[0043] Mechanical property characterization Bending strength was tested using the three-point bending method. The sintered sample was cut into strips of 3×4×20 mm. After beveling the strips onto a diamond plate, the samples were measured on a universal testing machine with a loading rate of 0.5 mm·min. -1 Record the load at which the test sample breaks. The calculation formula is as follows:

[0044] In the formula, The value is denoted as MPa for the bending strength of the test sample; F is the load at which the test sample breaks; L is the span of the bending mold for the test sample, which is 16 mm in this experiment; b is the width of the test sample, mm; and ℎ is the height of the sample, mm.

[0045] Example 1 This embodiment provides a method for preparing silicon carbide ceramics with both high thermal conductivity and flexural strength, including the following steps: S1. SiC powder and ternary sintering aid are mixed and loaded into a polyurethane ball mill jar. Polyurethane grinding balls are added at a ball-to-powder ratio (i.e., the ratio of the mass of the polyurethane grinding balls to the sum of the masses of SiC powder and ternary sintering aid) of 4:1. Anhydrous ethanol is added as a dispersion medium, with a mass ratio of anhydrous ethanol to powder of 1:1 (i.e., the ratio of the mass of ethanol to the sum of the masses of SiC powder and ternary sintering aid is 1:1). The mixture is ball-milled in a planetary ball mill at 150 r / min for 10 h to ensure thorough and uniform mixing of all components. After mixing, the polyurethane grinding balls are removed, and the resulting slurry is transferred to a glass dish. It is then dried in an 80 °C vacuum drying oven for 24 hours. h. Thoroughly remove anhydrous ethanol from the slurry, crush the dried mixed powder using an agate mortar, and pass it through an 80-mesh nylon sieve to obtain a mixed powder with uniform particle size and good flowability; the ternary sintering aid includes Re2O3 powder, ReF3 powder and Bi2O3 powder, where Re is La; the mass fraction of SiC powder in the mixed powder is 96.7%, the sum of the mass fractions of Re2O3 and ReF3 in the mixed powder is 3%, the mass fraction of Bi2O3 in the mixed powder is 0.3%, and the molar ratio of Re2O3 to ReF3 is 1:1; S2. The mixed powder from S1 is loaded into a graphite mold with an inner diameter of 25 mm. Before loading, graphite paper matching the inner wall size of the graphite mold is pre-cut and tightly attached to the inner wall of the mold and the end faces of the upper and lower pressure columns to prevent the sintered sample from sticking to the mold at high temperature and to facilitate demolding of the sample after sintering. The mold is placed in a spark plasma sintering furnace, and a uniaxial pressure of 40 MPa is applied under a vacuum atmosphere. The temperature is heated to 1900 °C and maintained at this temperature for 10 min. The sintering process is then ended. After the furnace temperature reaches room temperature, the mold is removed, and the sample is demolded using a tablet press. The graphite paper on the surface of the sample is removed using a grinding machine. The final sample with a diameter of approximately 25 mm is obtained, which is silicon carbide ceramic (denoted as 3LaF-BiO).

[0046] Furthermore, following the method described in Example 1 above, the composition of the rare earth sintering aid (the molar ratio of Re2O3 and ReF3 in the ternary sintering aid is fixed at 1:1) in the mixed powder is changed, and the remaining process parameters are the same as in Example 1. The prepared silicon carbide ceramics are denoted as 3YbF-BiO, 3ScF-BiO, and 3YF-BiO, as shown in Table 1 below.

[0047] Table 1 - SiC mass fraction and composition of ternary sintering aids in different mixed powders ; At a sintering temperature of 1900 °C, the densification behavior of SiC ceramics prepared by different rare earth element systems in Example 1 showed significant differences, as shown in Table 2.

[0048] Table 2. Densification and Properties of Bismuth Oxide Added with Different Rare Earth Oxides-Fluorides in Liquid Phase Sintering of Silicon Carbide Ceramics ; Due to the large ionic radius of rare earth element La, the 3LaF-BiO sample readily reacts with SiO2 to form a liquid phase. This liquid phase has low viscosity and is prone to volatilization during high-temperature sintering, resulting in insufficient actual liquid phase participation in densification. Consequently, SiC particles struggle to rearrange adequately through the liquid phase, leaving numerous pores trapped at grain boundaries. In contrast, Yb and Sc, with their smaller ionic radii, often exhibit higher viscosity liquid phases. While these liquid phases have slightly lower fluidity, they are more stable at high temperatures, less prone to volatilization, and can remain stably at grain boundaries, fully utilizing the liquid phase sintering effect. This promotes mass transfer and porosity reduction. Compared to SiC ceramics prepared with binary sintering aids without Bi2O3, the use of ternary sintering aids further improves the sintering density of the smaller ionic radii of the 3YbF-BiO and 3ScF-BiO samples.

[0049] Figure 1XRD diffraction patterns of SiC ceramics prepared with different rare earth element additive systems in Example 1 ( Figure 1 The image shows the XRD pattern of the SiC powder used as raw material. Analysis of the XRD diffraction patterns revealed that the main crystal forms of all SiC ceramic samples were 6H and 4H, while different rare earth element systems showed significant differences in the formation of secondary phases. Only BiF3 was detected in the 3LaF-BiO sample, while rare earth silicate phases were detected in the 3YF-BiO, 3YbF-BiO, and 3ScF-BiO samples. Rare earth silicates are formed by the reaction of sintering aids and SiO2 on the SiC surface. 3+ With the largest ionic radius, it forms the lowest viscosity liquid phase with SiO2, making it prone to volatilization and loss at high temperatures. The formation of BiF3 can be attributed to a fluorine-oxygen exchange reaction: during high-temperature sintering, Bi2O3 in the raw materials undergoes anion exchange with ReF3, with fluoride ions migrating from rare earth fluorides to the Bi-O system to form BiF3, while oxygen ions enter the rare earth silicate network. Furthermore, Sc... 3+ The ionic radius is the smallest and the liquid phase viscosity is high, making it difficult for rare earth silicates to form. Therefore, the second phase peak intensity of the 3ScF-BiO sample is the weakest.

[0050] The microstructure and grain size of SiC ceramics with different rare earth element additive systems in Example 1 are statistically analyzed as follows: Figure 2 As shown. Figure 2 In the middle, (a) and (e) are SEM images and grain size distributions of 3LaF-BiO, respectively; (b) and (f) are SEM images and grain size distributions of 3YF-BiO, respectively; (c) and (g) are SEM images and grain size distributions of 3YbF-BiO, respectively; and (d) and (h) are SEM images and grain size distributions of 3YbF-BiO, respectively.

[0051] Bi₂O₃ plays completely different roles in different rare earth systems, which is related to the fact that the ionic radius of rare earth elements determines the liquid phase viscosity and the thickness of the liquid phase layer. For the La system, the ionic radius is relatively large, resulting in a low viscosity, good fluidity, and a thin liquid phase layer when it forms with Re₂O₃-ReF₃-Bi₂O₃. According to liquid phase sintering theory, a low-viscosity liquid phase reduces grain boundary migration resistance, allowing grain growth to be controlled by interfacial reactions. This leads to a rapid dissolution and precipitation rate, and the grains can coarsen quickly. Under this mechanism, the average grain size (i.e., D) of 3LaF-BiO and 3YF-BiO is significantly different. 50The grain sizes reached 2.51 μm and 3.74 μm respectively, significantly larger than the sample without Bi₂O₃. For the Yb and Sc systems, the small ionic radius and strong cation field led to the formation of a high-viscosity liquid phase during sintering, with a significantly increased liquid phase thickness. According to the diffusion-controlled grain growth formula, the grain growth rate is inversely proportional to the liquid phase thickness. The thicker the liquid phase, the longer the diffusion path of the solute in the liquid phase, the greater the resistance, and the grain growth rate decreases significantly, transforming the growth mechanism into diffusion control. Therefore, the Dg of 3YbF-BiO and 3ScF-BiO is significantly higher. 50 The grain sizes are only 1.77 μm and 1.08 μm, and the growth rate is much smaller than that of the Y and La systems.

[0052] like Figure 3 As shown, to reveal the influence of Bi2O3 on the Yb-based sintering aid system samples, transmission electron microscopy was used to observe and analyze the grain boundary structure and second-phase composition of the 3YbF-BiO sample. Specifically, Figure 3 In the image, (a) is a STEM image of the 3YbF-BiO sample, (b) is a HAADF image, (c) is an HRTEM image, and (d) is a Fast Fourier Transform (FFT) image of the labeled region.

[0053] From the perspective of the second phase composition, after fast Fourier transform calibration of the HRTEM image of the 3YbF-BiO sample, the second phase was confirmed to be Yb2Si2O7. The addition of Bi2O3 promoted the reaction between YbF3 and Bi2O3 through fluorine-oxygen exchange, thereby consuming YbF3 and inhibiting the reaction between YbF3 and SiO2 to form Yb2O3, resulting in a more homogenized second phase. Furthermore, the thermal conductivity of Yb2Si2O7 is 4.31 W / (m·K), higher than that of Yb2O3 (2.1 W / (m·K)). Therefore, the contribution of single Yb2Si2O7 to phonon scattering is smaller compared to the mixed second phase, which is beneficial for heat transport and improves the thermal conductivity of the 3YbF-BiO sample.

[0054] from Figure 3 As shown in (b), the amorphous layer thickness between adjacent SiC grains in the 3YbF-BiO sample is only 1.11 nm. The amorphous film typically originates from the glassy phase formed during cooling by sintering aids, and its thickness is influenced by the amount of liquid phase and grain boundary wetting behavior. After adding Bi₂O₃, on the one hand, the fluorine-oxygen exchange reaction consumes some fluoride, altering the composition and viscosity of the liquid phase; on the other hand, Bi₂O₃, as a low-melting-point oxide, may promote the uniform spreading of the liquid phase at the grain boundaries, making it easier for the crystalline phase to precipitate from the grain boundary layer during cooling rather than remaining as a thick glassy layer. A thinner amorphous film means a reduction in the glassy phase at the grain boundaries, which is beneficial for phonon transport because the scattering of phonons by disordered glass is much stronger than that of a thin, well-crystallized interface. The thin amorphous film is another reason for the improved thermal conductivity.

[0055] The Re2O3-ReF3-Bi2O3 sample showed significant differences compared to the same system without Bi2O3. The thermal conductivity and phonon mean free path generally increased. This change is closely related to the regulatory effect of Bi2O3 on liquid phase behavior, phase composition, densification process, and grain growth. The thermal conductivity of different rare earth systems is also affected differently by Bi2O3. Samples with larger rare earth ionic radii achieve higher thermal conductivity, mainly due to their larger grain size. Coarse grains mean a relatively smaller number of grain boundaries per unit volume, and since grain boundaries are one of the main sources of phonon scattering, a reduction in their number is beneficial for phonon transmission. For the 3YbF-BiO sample with smaller rare earth ionic radii, its thermal conductivity decreased to 144.3 W / (m·K), which is the result of multiple factors. First, the grain size is small, only 1.77 μm, resulting in a significant increase in the number of grain boundaries, which significantly enhances phonon scattering at the grain boundaries. Secondly, the uniform and continuous distribution of the second phase around the SiC grains complicates the grain boundary region, causing phonons to encounter more heterogeneous interfaces when crossing grain boundaries, thus increasing the scattering probability. Although the sample has extremely low porosity, eliminating the adverse effects of pore scattering, the scattering effect of fine grains and the second phase is more prominent, leading to a decrease in the phonon mean free path to 16.6 nm and a further reduction in thermal conductivity. The 3ScF-BiO sample, with the smallest rare-earth ion radius, also has the lowest thermal conductivity. This is closely related to the extremely fine grain size of the Sc system; 1.08 μm grains mean the highest grain boundary density, resulting in frequent grain boundary scattering of phonons during transmission. The combined effect of extremely fine grains and the high-scattering-intensity grain boundary phase significantly reduces the phonon mean free path to 12.8 nm. Although this system has the highest density, its thermal conductivity remains the lowest.

[0056] The flexural strength initially increases and then decreases with decreasing rare earth ion radius; the 3LaF-BiO sample exhibits the lowest strength at 340 MPa, which is directly related to its coarse microstructure. SEM images reveal numerous pores in the 3LaF-BiO sample. These pores act as stress concentration sources, inducing localized stress amplification under load and becoming the starting point for crack initiation. The flexural strength of the 3YF-BiO sample is significantly higher than that of the 3LaF-BiO sample, reaching 453 MPa. Although it has the largest grain size, and according to the classic Hall-Page relation, coarser grains generally indicate lower strength, the 3YF-BiO sample achieves higher strength, demonstrating that grain size is not the sole determinant of strength. Furthermore, its porosity is much lower than that of the 3LaF-BiO sample, making the stress concentration effect relatively controllable. The 3YbF-BiO sample exhibits the highest strength, reaching 463 MPa. This is closely related to its excellent microstructure, with a density as high as 98.8% and extremely low porosity, fundamentally eliminating the adverse effects of pores as crack initiation sites. Simultaneously, this sample has a grain size of 1.77 μm, which can enhance strength through a grain refinement strengthening mechanism. Smaller grains mean a larger grain boundary area per unit volume, requiring crack propagation to constantly change direction or consume more energy. The 3ScF-BiO sample has the highest density, almost eliminating the adverse effects of pores, and the finest grain size at only 1.08 μm. From the perspective of grain refinement strengthening, it should theoretically achieve the highest strength, but its actual strength is lower than that of the Yb system. This seemingly anomalous phenomenon needs to be explained by the amount of residual second phase. Although rare earth silicate phases can act as fillers in the intergranular pores, indirectly improving the density of silicon carbide ceramics, the presence of this intergranular phase leads to a decrease in the intergranular bonding strength of silicon carbide grains due to the non-in-situ heterogeneous interface between this second phase and the silicon carbide grains. This indirectly results in a decrease in the flexural strength of the sample. Furthermore, since the flexural strength of rare earth silicate phases is generally lower than that of silicon carbide, the higher the amount of residual rare earth silicate phase, the more significant the decrease in the flexural strength of the silicon carbide ceramic. Therefore, although the 3YF-BiO and 3YbF-BiO samples exhibit excellent sintering performance, the high liquid phase viscosity and large amount of residual second phase mean that the negative impact of the intergranular phase on the intergranular bonding strength outweighs the strengthening effect brought about by the increase in sample density.

[0057] Comparative Example 1 This comparative example provides a method for preparing silicon carbide ceramics, including the following steps: S1. SiC powder and binary sintering aid are mixed and loaded into a polyurethane ball mill jar. Polyurethane grinding balls are added at a ball-to-powder ratio (i.e., the ratio of the mass of polyurethane grinding balls to the sum of the masses of SiC powder and ternary sintering aid) of 4:1. Anhydrous ethanol is added as a dispersion medium, with a mass ratio of anhydrous ethanol to powder of 1:1 (i.e., the ratio of the mass of ethanol to the sum of the masses of SiC powder and ternary sintering aid is 1:1). The mixture is ball-milled in a planetary ball mill at 150 r / min for 10 h to ensure thorough and uniform mixing of all components. After mixing, the polyurethane grinding balls are removed, and the resulting slurry is transferred to a glass dish. It is then dried in an 80 °C vacuum drying oven for 24 hours. h. Thoroughly remove anhydrous ethanol from the slurry, crush the dried mixed powder using an agate mortar, and pass it through an 80-mesh nylon sieve to obtain a mixed powder with uniform particle size and good flowability; the binary sintering aid includes Re2O3 powder and ReF3 powder, where Re is La; the mass fraction of SiC powder in the mixed powder is 97%, the sum of the mass fractions of Re2O3 and ReF3 in the mixed powder is 3%, and the molar ratio of Re2O3 to ReF3 is 1:1; S2. The mixed powder from S1 is loaded into a graphite mold with an inner diameter of 25 mm. Before loading the mold, graphite paper that matches the inner wall size of the graphite mold is pre-cut and tightly attached to the inner wall of the mold and the end faces of the upper and lower pressure columns to prevent the sintered sample from sticking to the mold at high temperature and to facilitate demolding of the sample after sintering. The mold is placed in a spark plasma sintering furnace, and a uniaxial pressure of 40 MPa is applied under a vacuum atmosphere. The temperature is heated to 1900 °C and maintained at this temperature for 10 min. The sintering process is then ended. After the furnace temperature reaches room temperature, the mold is removed, and the sample is demolded using a tablet press. The graphite paper on the surface of the sample is removed using a grinding machine. The final sample with a diameter of approximately 25 mm is obtained, which is silicon carbide ceramic (denoted as 3LaF).

[0058] Furthermore, following the method in Comparative Example 1 above, the types of rare earth element Re in the binary sintering aid were changed to Yb, Sc, and Y (the molar ratio of Re2O3 and ReF3 in the sintering aid was fixed at 1:1), and the remaining process parameters were the same as in Comparative Example 1. The prepared silicon carbide ceramics were denoted as 3YbF, 3ScF, and 3YF, as shown in Table 3 below.

[0059] Table 3. SiC mass fraction and composition of binary sintering aids in different mixed powders. ; The densification results of SiC ceramics prepared by different rare earth element systems in Comparative Example 1 are shown in Table 4. Compared with Example 1, their densification and porosity show similar changes.

[0060] Table 4 - Densification and Properties of Silicon Carbide Ceramics Sintered from Different Rare Earth Oxides and Fluorides in Liquid Phase ; As shown in Table 4, the thermal conductivity of the 3ScF and 3YbF samples is 124 W / (m·K) and 123 W / (m·K), respectively. The 3LaF sample, due to insufficient densification, has a moderate thermal conductivity. The 3YF sample, with the lowest degree of densification, more porosity and defects, and smaller grain size, has the lowest thermal conductivity among the four, at only 102 W / (m·K). Except for Sc, all samples have lower thermal conductivity than the sample with added Bi₂O₃ in Example 1 (highest being 170 W / (m·K)). A comprehensive comparison of the two groups of samples revealed two positive effects from the addition of Bi2O3: First, the second phase transformed from a mixture of RE2Si2O7 and RE2O3 into a single RE2Si2O7 phase, eliminating the low thermal conductivity of RE2O3 and resulting in a more uniform grain boundary composition. The thinner intergranular amorphous film reduced the scattering of phonons by the glass. These microstructural optimizations collectively improved the thermal conductivity of the samples. Second, except for La, the closed porosity was significantly reduced, improving the flexural strength of the samples.

[0061] Example 2 This embodiment provides a method for preparing silicon carbide ceramics with both high thermal conductivity and flexural strength, including the following steps: S1. SiC powder and ternary sintering aid are mixed and loaded into a polyurethane ball mill jar. Polyurethane grinding balls are added at a ball-to-powder ratio (i.e., the ratio of the mass of the polyurethane grinding balls to the sum of the masses of SiC powder and ternary sintering aid) of 4:1. Anhydrous ethanol is added as a dispersion medium, with a mass ratio of anhydrous ethanol to powder of 1:1 (i.e., the ratio of the mass of ethanol to the sum of the masses of SiC powder and ternary sintering aid is 1:1). The mixture is ball-milled in a planetary ball mill at 150 r / min for 10 h to ensure thorough and uniform mixing of all components. After mixing, the polyurethane grinding balls are removed, and the resulting slurry is transferred to a glass dish. It is then dried in an 80 °C vacuum drying oven for 24 hours. h. Thoroughly remove anhydrous ethanol from the slurry, crush the dried mixed powder using an agate mortar, and pass it through an 80-mesh nylon sieve to obtain a mixed powder with uniform particle size and good flowability; the ternary sintering aid includes Y2O3 powder, YF3 powder and Bi2O3 powder; the mass fraction of SiC powder in the mixed powder is 96.7%, the sum of the mass fractions of Y2O3 and YF3 in the mixed powder is 3%, the mass fraction of Bi2O3 in the mixed powder is 0.3%, and the molar ratio of Re2O3 to ReF3 is 1:1; S2. The mixed powder from S1 is loaded into a graphite mold with an inner diameter of 25 mm. Before loading, graphite paper matching the inner wall size of the graphite mold is pre-cut and tightly attached to the inner wall of the mold and the end faces of the upper and lower pressure columns to prevent the sintered sample from sticking to the mold at high temperature and to facilitate demolding of the sample after sintering. The mold is placed in a spark plasma sintering furnace, and a uniaxial pressure of 40 MPa is applied under a vacuum atmosphere. The temperature is heated to 1900 °C and maintained at this temperature for 10 min. The sintering process is then ended. After the furnace temperature reaches room temperature, the mold is removed, and the sample is demolded using a tablet press. The graphite paper on the surface of the sample is removed using a grinding machine. The final sample with a diameter of approximately 25 mm is obtained, which is silicon carbide ceramic (denoted as 3YF-BiO).

[0062] Furthermore, following the method in Example 2 above, the composition of rare earth sintering aid (the molar ratio of Y2O3 and YF3 in the ternary sintering aid is fixed at 1:1) in the mixed powder is changed, and the remaining process parameters are the same as in Example 2. The prepared silicon carbide ceramics are denoted as 5YF-BiO and 7YF-BiO, as shown in Table 5 below.

[0063] Table 5 - SiC mass fraction and composition of ternary sintering aids in different mixed powders ; The sintering properties of SiC ceramics with different contents (3 wt.%, 5 wt.%, 7 wt.%) using Y2O3-YF3-Bi2O3 as a sintering aid in Example 2 are shown in Table 6.

[0064] Table 6 - Densification Results of Silicon Carbide Ceramics with Different Liquid Phase Contents ; As shown in Table 6, the densification of the samples initially increased and then decreased as the additive content increased from 3 wt.% to 7 wt.%. The 3YF-BiO sample exhibited lower densification, which is closely related to insufficient liquid phase during liquid-phase sintering. When the additive content was only 3 wt.%, the total amount of liquid phase formed at high temperature was limited, making it difficult to fully fill the gaps between SiC particles. In the early stages of sintering, the liquid phase needed to rely on capillary force to drive particle rearrangement, but insufficient liquid phase led to incomplete particle rearrangement. After entering the dissolution-precipitation stage, the limited liquid phase could not provide sufficient mass transfer channels for the dissolution-transport process of SiC, and a large number of pores were retained inside the sample and could not be discharged, resulting in a density of only 91.4%. The 5YF-BiO sample showed the best densification effect, which was attributed to the appropriate amount of liquid phase fully utilizing the role of the sintering aid. When the additive content was increased to 5 wt.%, the amount of liquid phase generated at high temperature was sufficient to provide a stable medium for dissolution-precipitation mass transfer, achieving a densification level close to the theoretical density. When the content of additives was further increased to 7 wt.%, although the initial amount of liquid phase was greater, the excessive liquid phase was prone to volatilization at high temperature, resulting in a reduction in the actual amount of liquid phase participating in densification. Therefore, the porosity of the 7YF-BiO sample increased to 2.44%, and the density also decreased to 94.7%.

[0065] The microstructure of silicon carbide ceramics prepared in Example 2 using Y2O3-YF3-Bi2O3 as a sintering aid is as follows: Figure 4 As shown; Figure 4 In the image, (a) and (d) represent SEM images and grain size distribution of 3YF-BiO, (b) and (e) represent SEM images and grain size distribution of 5YF-BiO, and (c) and (f) represent SEM images and grain size distribution of 7YF-BiO.

[0066] from Figure 4 As shown in (a) to (c), as the content of the additives increased from 3 wt.% to 7 wt.%, the density of the microstructure first improved significantly and then slightly deteriorated. The 3YF-BiO sample contained a large number of pores, the 5YF-BiO sample showed almost no pores and exhibited a highly dense microstructure, and the 7YF-BiO sample, although containing fewer pores, still showed clear signs of pores. The second phase in all three samples was in an isolated and dispersed state.

[0067] like Figure 4 As shown in (d) to (f), the grain size initially increases and then decreases with increasing additive content. The D in the 3YF-BiO sample... 50The grain size was 3.74 μm. When the additive content was only 3 wt.%, the amount of liquid phase generated at high temperature was limited. The presence of pores hindered the long-range diffusion of the material, and the final grain size was much smaller than that of the 5YF-BiO sample (4.51 μm). As the additive content further increased, the viscosity of the excess liquid phase increased, which inhibited the mass transfer rate and the final grain size decreased.

[0068] With increasing additive content, the thermal conductivity of the samples initially increased and then decreased. The thermal conductivity of the 3YF-BiO sample was 170.7 W / (m·K), increasing to 197.1 W / (m·K) for the 5YF-BiO sample, while decreasing to 160.2 W / (m·K) for the 7YF-BiO sample. The 3YF-BiO sample exhibited numerous pores and a small grain size. The 5YF-BiO sample had a porosity of only 0.12%, almost eliminating the phonon scattering effect of pores, while its grain size increased to 4.51 μm, achieving the highest thermal conductivity. As the additive content increased to 7 wt.%, the thermal conductivity of the sample actually decreased, a result of the combined effect of decreasing grain size and increasing porosity.

[0069] As the additive content increased from 3 wt.% to 7 wt.%, the flexural strength showed a trend of first decreasing and then increasing, with the 3YF-BiO sample reaching a flexural strength of 453 MPa. The flexural strength of the 5YF-BiO sample reached 416 MPa, which was slightly lower due to the larger grain size. In the 7YF-BiO sample, the excessive liquid phase volatilized severely at high temperatures, resulting in a large number of residual pores, causing the flexural strength to drop significantly to 355 MPa.

[0070] In summary, this invention employs spark plasma sintering technology, using SiC powder as raw material and a Re2O3-ReF3-Bi2O3 (Re = La, Y, Yb, Sc) ternary system as sintering aids to prepare densified silicon carbide ceramics. The effects of rare earth element types (La, Y, Yb, Sc) and aid content (3~7 wt.%) on the microstructure, phase composition, and thermal and mechanical properties of liquid-phase sintered silicon carbide ceramics were systematically studied. Specific conclusions are as follows: 1. The type of rare earth element has a significant impact on the densification behavior, phase composition, and microstructure of SiC ceramics. The Yb and Sc system, due to its small ionic radius, high liquid phase viscosity, and good high-temperature stability, achieves high densification. In contrast, the La and Y system, due to its low liquid phase viscosity and significant volatilization, exhibits lower densification. The occurrence of the fluorine-oxygen exchange reaction suppresses the adverse effects of additive volatilization. The Y and Yb systems simultaneously generate rare earth silicate phases. Further research using Y-based additives reveals that the content of Y-based additives exhibits a typical optimal ratio for the densification behavior of SiC ceramics. At 3 wt.%, insufficient liquid phase leads to incomplete densification. At 5 wt.%, the liquid phase content is moderate, achieving complete densification with a density of 99.1% and an open porosity of 0.12%. However, when the additive content increases to 7 wt.%, excessive liquid phase leads to volatilization and a decrease in density.

[0071] 2. The type of rare earth elements significantly affects the thermal properties of SiC ceramics by regulating liquid phase characteristics and grain growth behavior. The Y-based system, with the largest grain size and the fewest grain boundaries, achieves the highest thermal conductivity, 170.7 W / (m·K); grain size and grain boundary characteristics are the core factors determining thermal conductivity. Studies on the content of Y-based additives revealed that grain size and porosity significantly affect thermal conductivity. The 5YF-BiO sample, due to its coarse grains and extremely low porosity, achieves a thermal conductivity of 197.1 W / (m·K). Although the 3YF-BiO sample has relatively large grains, the scattering effect of numerous pores limits its thermal conductivity to 170.7 W / (m·K). The 7 wt.% sample, due to grain refinement and increased porosity, has a thermal conductivity that decreases to 160.2 W / (m·K).

[0072] 3. The difference in strength is the result of the synergistic effect of porosity, grain size, and grain boundary bonding: the Yb system achieves fine-grain strengthening through high density, fine-grained structure, and uniformly distributed grain boundary phases; although the strength of the Y system is slightly lower than that of the Yb system, the strong bonding of grain boundaries still offsets the adverse effects of coarse grains and porosity; although the Sc system is extremely dense and has the finest grains, the presence of residual second phase weakens grain boundary bonding, limiting further improvement in strength; the La system suffers from strength degradation due to stress concentration dominated by a large number of pores. The effect of Y-based additive content on performance shows that although the 3YF-BiO sample has high porosity, its grain boundaries are clean and the bonding is strong; although the 7YF-BiO sample has the smallest grain size, its strengthening effect is weakened due to a large number of pores, resulting in the lowest flexural strength; the 5YF-BiO sample has an appropriate liquid phase content, complete densification, and a fine interface structure, which optimizes thermal properties while also taking into account flexural strength (416 MPa); overall, the 5YF-BiO sample has better comprehensive performance.

[0073] It is understood that the technical features of the above embodiments can be combined arbitrarily. For the sake of brevity, not all possible combinations of the technical features in the above embodiments are described. However, as long as there is no contradiction in the combination of these technical features, they should be considered to be within the scope of this specification.

[0074] The above are merely preferred embodiments of this application, and only specifically describe the technical principles of this application. These descriptions are only for explaining the principles of this application and should not be construed as limiting the scope of protection of this application in any way. Based on this explanation, any modifications, equivalent substitutions, and improvements made within the spirit and principles of this application, as well as other specific embodiments of this application that can be conceived by those skilled in the art without creative effort, should be included within the scope of protection of this application.

Claims

1. A method for preparing silicon carbide ceramics possessing both high thermal conductivity and flexural strength, characterized in that, Includes the following steps: SiC powder and ternary sintering aid are mixed to obtain a mixed powder. The mixed powder was subjected to spark plasma sintering to obtain silicon carbide ceramics; The ternary sintering aids include Re2O3, ReF3 and Bi2O3, wherein Re is selected from any one of La, Y, Yb and Sc.

2. The method for preparing silicon carbide ceramics as described in claim 1, characterized in that, The mass fraction of SiC powder in the mixed powder is 92.7-96.7%, the sum of the mass fractions of Re2O3 and ReF3 in the mixed powder is 3-7%, and the remainder is Bi2O3.

3. The method for preparing silicon carbide ceramics as described in claim 1, characterized in that, The molar ratio of Re2O3 to ReF3 is (1~2):(1~2).

4. The method for preparing silicon carbide ceramics as described in claim 1, characterized in that, The mixed powder is placed in a spark plasma sintering furnace, where a pressure of 38-40 MPa is applied under a vacuum atmosphere, and the mixture is heated to 1850-1950°C and held at this temperature for 10-15 minutes to obtain silicon carbide ceramics.

5. The method for preparing silicon carbide ceramics as described in claim 1, characterized in that, SiC powder, ternary sintering aid After mixing, ethanol is added, followed by ball milling and drying to obtain a mixed powder.

6. The method for preparing silicon carbide ceramics as described in claim 5, characterized in that, SiC powder, ternary sintering aid After mixing, ethanol is added and the mixture is placed in a ball mill jar. Grinding balls are added, and the mixture is ball-milled at 150-200 r / min for 10-15 h. After drying, the mixed powder is obtained.

7. The method for preparing silicon carbide ceramics as described in claim 6, characterized in that, The ratio of the mass of the grinding ball to the sum of the masses of SiC powder and ternary sintering aid is (4~6):

1.

8. The method for preparing silicon carbide ceramics as described in claim 6, characterized in that, The ratio of the mass of ethanol to the sum of the masses of SiC powder and ternary sintering aid is (1~2):(1~2).

9. The method for preparing silicon carbide ceramics as described in claim 6, characterized in that, SiC powder, ternary sintering aid After mixing, ethanol is added and placed in a ball mill jar. Grinding balls are added, and the mixture is ball-milled at 150-200 r / min for 10-15 h. The mixture is then dried at 80-100℃ for 20-24 h, pulverized, and passed through an 80-100 mesh sieve to obtain the mixed powder.

10. A silicon carbide ceramic possessing both high thermal conductivity and flexural strength, characterized in that, It is prepared by any one of the preparation methods described in claims 1 to 9.

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