A multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic and its preparation method

Abstract

A multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic and a preparation method thereof, which belongs to the field of controllable preparation of ultra-high temperature ceramics. The purpose of the present invention is to solve the problem that the sintering temperature of the current preparation of multi-element ultra-high temperature solid solution ceramics is high, low-temperature densification is difficult, and ceramics with excellent performance cannot be prepared under low temperature conditions. The present invention uses HfSi2, TaSi2, B4C and C as raw materials, and is prepared by in-situ reaction combined with rapid hot pressing sintering at low temperature; densification is achieved in only 5 to 15 minutes at a temperature as low as 1600 to 1800°C; the (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic prepared by the present invention exhibits better mechanical properties than HfB2-HfC-SiC ceramics, and the introduction of Ta elements easily forms a low melting point phase and helps to form a dense continuous ablation layer, which can significantly improve the ablation resistance.

Description

Technical Field

[0001] The present invention belongs to the field of controllable preparation of ultra-high temperature ceramics, and specifically relates to a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic and a preparation method thereof. Background Art

[0002] With the development of aerospace technology, the Mach number of missiles and various aircraft has gradually increased. The thermal protection materials of aircraft hot end components are facing the threat of harsh service environment such as high-speed airflow particle scouring, oxidation and gas corrosion. The research and development and design of thermal protection structural materials have ushered in new requirements and challenges.

[0003] Ultra-high temperature ceramics are candidate materials for high temperature thermal structure applications under extraordinary environments due to their excellent heat resistance, oxidation resistance and thermochemical stability. Hf-based ultra-high temperature ceramics (HfB2 and HfC, etc.) have high melting point, high hardness, excellent heat resistance, oxidation resistance and thermochemical stability, and are widely used in the aerospace industry, including the wing leading edge, end cap and hot end of the engine of missiles and space shuttles. However, it is difficult to obtain dense bulk pure HfB2 ceramics by sintering, which is attributed to its strong covalent bond binding energy and low self-diffusion coefficient (Journal of the European Ceramic Society, 2018, 28 (6): 1287-1296.). Studies have shown that the introduction of SiC can promote the densification of HfB2 ceramics and improve mechanical properties (Journal of the European Ceramic Society, 2012, 32 (10): 2557-2563; Composites Science and Technology, 2005, 65 (12): 1869-1879.). On this basis, the researchers found that HfB2-HfC-SiC ternary ceramics with HfC showed more uniform microstructure and excellent mechanical properties. However, the current HfB2-HfC-SiC ternary ceramics have high preparation temperature, long sintering time, a small amount of intergranular fracture (the main fracture mode is transgranular fracture) and small-sized grains pulled out (with high orientation), which makes the mechanical properties of ternary ceramics poor. Introducing Ta elements into hafnium-based ultra-high temperature ceramics can effectively promote densification. Ta x Hf 1-xC solid solution ceramics exhibit excellent properties of high strength, oxidation resistance and ablation resistance in extremely high temperature environments (Acta Materialia, 2018, 145: 142-153; Corrosion Science, 2020, 164: 108348; Ceramics International, 2017, 43 (17): 14798-14806.). Studies have shown that the introduction of Ta elements into HfC-SiC ceramics to form Hf x Ta 1-x C-SiC solid solution ceramics can significantly improve the mechanical properties of HfC-SiC ceramics and obtain a uniformly distributed microstructure. The obvious intergranular fracture and the extraction of large-sized grains make Hf x Ta 1-x The anti-destructive ability of C-SiC solid solution ceramics is enhanced (Journal of the European Ceramic Society, 2018, 38 (4): 1227-1236.). There are also studies reporting that the introduction of Ta elements can promote the densification of HfB2-SiC ceramics. The prepared (Hf, Ta) B2-SiC solid solution ceramics have a fine microstructure. However, the fracture toughness of (Hf, Ta) B2-SiC ceramics is lower than that of HfB2-SiC ceramics. The reason is that as the content of TaB2 solid solution powder increases, the cross section of the sample becomes smoother, and the transgranular fracture mode plays a leading role (Journal of Synthetic Crystals, 2019, 48 (12): 2303-2307.). In summary, the sintering temperature for preparing multi-element ultra-high temperature solid solution ceramics is high, and low-temperature densification is difficult. How to prepare dense multi-component ultra-high temperature boride-carbide-silicon carbide ceramics at a lower temperature so that they have the excellent properties of HfB2-HfC-SiC ternary ceramics and Hf-Ta solid solution ceramics is a problem that needs to be solved urgently.

[0004] In the prior art, the transition metal Ta is introduced into Hf-based ultra-high temperature ceramics to improve the high temperature stability of Hf-based ultra-high temperature ceramics in extreme environments. The specific methods are: 1) The Ta element is introduced into HfC-SiC ceramics by mechanical alloying combined with pressureless sintering process to obtain (Hf, Ta) C-SiC solid solution ceramics, but its sintering temperature is high and low-temperature densification is difficult (Journal of the European Ceramic Society, 2018, 38 (4): 1227-1236.); 2) TaB2 is introduced into HfB2-SiC ceramics by boron thermal reduction combined with spark plasma sintering to prepare (Hf, Ta) B2-SiC solid solution ceramics, but its hardness and fracture toughness are low (Journal of Synthetic Crystals, 2019, 48 (12): 2303-2307.). Summary of the invention

[0005] The purpose of the present invention is to solve the problems of high sintering temperature, difficulty in low-temperature densification, and inability to prepare multi-component ultra-high temperature ceramics with excellent performance under low-temperature conditions in the current preparation of multi-component ultra-high temperature solid solution ceramics, and to provide a multi-component ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic and its preparation method.

[0006] The invention uses HfSi2, TaSi2, B4C and C as raw materials, and is prepared by in-situ reaction combined with rapid hot pressing and sintering at low temperature; densification is achieved in only 5min to 15min at a temperature as low as 1600℃ to 1800℃; the multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic has a flexural strength of 404.69±33.78MPa to 556.61±24.77MPa, and a fracture toughness of 3.23±0.12MPa·m 1 / 2 ~5.13±0.18MPa·m 1 / 2 , and has good ablation resistance at 2100K-2300K. The present invention adopts in-situ reaction and solid solution strengthening to make (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics show better mechanical properties than HfB2-HfC-SiC ternary ceramics.

[0007] A method for preparing multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics is specifically completed by the following steps:

[0008] 1. Preparation of composite powder with fine grains and uniform distribution by low-energy ball milling:

[0009] First, ZrO2 balls are added to a ball mill, and then HfSi2, TaSi2, B4C, C and anhydrous ethanol are poured into the ball mill, and then the ball mill is sealed; the sealed ball mill is fixed on a planetary ball mill, and argon is filled into the ball mill while being evacuated, so that the ball mill is subjected to low-energy ball milling for a period of time in an argon environment to obtain a ball-milled composite powder; the ball-milled composite powder is placed in a vacuum drying oven and dried at room temperature, and then sieved through a 120-200 mesh sieve in a glove box filled with argon to obtain a composite powder with refined grains and uniform distribution, and then stored in a vacuum;

[0010] The molar ratio of the total amount of HfSi2 and TaSi2 to B4C and C in step 1 is (1-6):(1-1.2):(3.6-18);

[0011] The molar ratio of HfSi2 to TaSi2 described in step 1 is (0.01-0.99):(0.01-0.99);

[0012] 2. Manually pre-pressing the composite powder stored in vacuum by a hydraulic press to obtain a composite powder green body;

[0013] 3. First, the temperature of the composite powder green body is raised from room temperature to 700°C; then, the axial pressure of the pressure head is raised to 30MPa~60MPa at a pressure increase rate of 1T / min~1.3T / min, and during the pressure increase process, the temperature of the composite powder green body is raised from 700°C to 1600°C~1800°C at a heating rate of 76.3°C / min~85.5°C / min, and then hot-pressed and sintered at 1600°C~1800°C and an axial pressure of 30MPa~60MPa for 5min~15min to obtain multi-element ultra-high temperature (Hf,Ta)B2-(Hf,Ta)C-SiC ceramics.

[0014] Beneficial effects of the present invention:

[0015] 1. The present invention adopts in-situ reaction combined with rapid hot pressing sintering process to realize low-temperature rapid preparation of multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics, effectively solving the key problems of high sintering temperature and difficult densification of multi-element ultra-high temperature ceramics;

[0016] Second, the in-situ reaction and solid solution behavior can effectively promote the densification process of multiphase ceramics. An interlocking structure of high aspect ratio grains is formed in the ceramic matrix, and the mechanical properties are significantly improved. The pull-out, crack deflection and bridging of high aspect ratio (Hf, Ta) B2 grains with interlocking structure are the main toughening mechanisms.

[0017] 3. The introduction of Ta element can easily form a low melting point phase and help form a dense continuous ablation layer, which can significantly improve the ablation resistance. When Hf / Ta=0.9:0.1, the best ablation resistance is shown, which is attributed to the Ta2O5 and Hf6Ta2O generated by ablation. 17 The oxides have a good bonding effect, forming a continuous and dense ablation protection layer, which slows down the erosion of the ablation layer by air flow scouring and protects the internal ceramic matrix from continuous damage. BRIEF DESCRIPTION OF THE DRAWINGS

[0018] Figure 1 The present invention is a flow chart for preparing the (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics;

[0019] Figure 2 : XRD diagram of the composite powder after ball milling in Examples 1 to 3 and Comparative Examples 1 to 2, in which Hf / Ta=0:1 is Comparative Example 2, Hf / Ta=0.5:0.5 is Example 3, Hf / Ta=0.7:0.3 is Example 2, Hf / Ta=0.9:0.1 is Example 1, and Hf / Ta=1:0 is Comparative Example 1;

[0020] Figure 3 The densification curve of H7T3B1C1-S ceramic prepared under the sintering conditions of 1700°C / 40MPa / 15min in Example 2;

[0021] Figure 4 XRD patterns of (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics with different Hf / Ta molar ratios prepared under sintering conditions of 1700°C / 40 MPa / 15 min in Examples 1-3 and Comparative Examples 1-2;

[0022] Figure 5 BSE diagrams of (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics with different Hf / Ta molar ratios prepared under sintering conditions of 1700°C / 40 MPa / 15 min in Examples 1 to 3 and Comparative Examples 1 to 2, wherein (a1) and (a2) are H10T0B1C1-S, (b1) and (b2) are H9T1B1C1-S, (c1) and (c2) are H7T3B1C1-S, (d1) and (d2) are H5T5B1C1-S, and (e1) and (e2) are H0T10B1C1-S;

[0023] Figure 6(a) is a TEM image of the H7T3B1C1-S ceramic prepared under the sintering conditions of 1700°C / 40MPa / 15min in Example 2; (b) is a HRTEM image of the H7T3B1C1-S ceramic prepared under the sintering conditions of 1700°C / 40MPa / 15min in Example 2;

[0024] Figure 7 (a)-(b) are dark field images and HADDF images of H7T3B1C1-S ceramics prepared under the sintering conditions of 1700°C / 40MPa / 15min in Example 2; (c)-(h) are element distribution diagrams of H7T3B1C1-S ceramics prepared under the sintering conditions of 1700°C / 40MPa / 15min in Example 2;

[0025] Figure 8 The fracture morphologies of (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics with different Hf / Ta molar ratios prepared under sintering conditions of 1700°C / 40 MPa / 15 min in Examples 1 to 3 and Comparative Examples 1 to 2, wherein (a) is H10T0B1C1-S, (b) is H9T1B1C1-S, (c) is H7T3B1C1-S, (d) is H5T5B1C1-S, and (e) is H0T10B1C1-S;

[0026] Fig. 9 XRD patterns of (Hf,Ta)B2-(Hf,Ta)C-SiC ceramics with different B4C / C molar ratios prepared in Examples 4 to 7 under the sintering conditions of 1700°C / 40MPa / 15min. DETAILED DESCRIPTION

[0027] Specific implementation method 1: The ceramic described in this implementation method is made of HfSi2, TaSi2, B4C and C as raw materials, and is quickly prepared by in-situ reaction combined with rapid hot pressing and sintering at low temperature; densification is achieved in only 5min to 15min at a temperature as low as 1600℃ to 1800℃; the bending strength of the multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic is 404.69±33.78MPa~556.61±24.77MPa, and the fracture toughness is 3.23±0.12MPa·m 1 / 2 ~5.13±0.18MPa·m 1 / 2 , has good ablation resistance at 2100K ~ 2300K.

[0028] Specific embodiment 2: This embodiment differs from the specific embodiment 1 in that: a method for preparing multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics is specifically completed in the following steps:

[0029] 1. Preparation of composite powder with fine grains and uniform distribution by low-energy ball milling:

[0030] First, ZrO2 balls are added to a ball mill, and then HfSi2, TaSi2, B4C, C and anhydrous ethanol are poured into the ball mill, and then the ball mill is sealed; the sealed ball mill is fixed on a planetary ball mill, and argon is filled into the ball mill while being evacuated, so that the ball mill is subjected to low-energy ball milling for a period of time in an argon environment to obtain a ball-milled composite powder; the ball-milled composite powder is placed in a vacuum drying oven and dried at room temperature, and then sieved through a 120-200 mesh sieve in a glove box filled with argon to obtain a composite powder with refined grains and uniform distribution, and then stored in a vacuum;

[0031] The molar ratio of the total amount of HfSi2 and TaSi2 to B4C and C in step 1 is (1-6):(1-1.2):(3.6-18);

[0032] The molar ratio of HfSi2 to TaSi2 described in step 1 is (0.01-0.99):(0.01-0.99);

[0033] 2. Manually pre-pressing the composite powder stored in vacuum by a hydraulic press to obtain a composite powder green body;

[0034] 3. First, the composite powder green body is heated from room temperature to 700°C; then the axial pressure of the pressure head is increased to 30MPa-60MPa at a pressure increase rate of 1T / min-1.3T / min, and the temperature of the composite powder green body is increased from 700°C to 1600°C-1800°C at a heating rate of 76.3°C / min-85.5°C / min during the pressure increase process, and then hot-pressed and sintered for 5min-15min at 1600°C-1800°C and an axial pressure of 30MPa-60MPa to obtain multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics. The other steps are the same as those in the first embodiment.

[0035] Specific implementation method 3: This implementation method is different from specific implementation method 1 or 2 in that: the ZrO2 balls described in step 1 are large balls and small balls, the diameter of the large balls is 9.95mm-10.05mm, the diameter of the small balls is 5.95mm-6.05mm, and the mass ratio of the large balls to the small balls is (3-5):(1-2). The other steps are the same as specific implementation method 1 or 2.

[0036] Specific embodiment 4: This embodiment differs from specific embodiments 1 to 3 in that the total mass ratio of the ZrO2 balls described in step 1 to the total mass ratio of HfSi2, TaSi2, B4C and C is (6-15): 1. The other steps are the same as specific embodiments 1 to 3.

[0037] Specific embodiment 5: This embodiment differs from specific embodiments 1 to 4 in that the anhydrous ethanol described in step 1 accounts for 31.6% to 47.4% of the total mass of HfSi2, TaSi2, B4C and C. The other steps are the same as those of specific embodiments 1 to 4.

[0038] Specific implementation method 6: This implementation method is different from specific implementation methods 1 to 5 in that: the rotation speed of the low-energy ball mill described in step 1 is 250r / min to 350r / min, the time of the low-energy ball mill is 24h to 48h; the time of room temperature drying described in step 1 is 24h to 36h. The other steps are the same as specific implementation methods 1 to 5.

[0039] Specific embodiment 7: The difference between this embodiment and specific embodiments 1 to 6 is that: in step 2, the graphite paper coated with a release agent is closely attached to the wall of the graphite mold, the composite powder stored in vacuum is poured into the circular graphite mold, and then the graphite gasket and the composite powder, and the composite powder and the pressure head are separated by graphite paper coated with a release agent; the graphite mold containing the composite powder is placed under a hydraulic press and manually pressed to 3MPa to 5MPa to obtain a composite powder embryo. The other steps are the same as specific embodiments 1 to 6.

[0040] Specific embodiment 8: This embodiment differs from specific embodiments 1 to 7 in that: the release agent is one or more of talcum powder, clay and white clay; the diameter of the graphite mold is 30 mm to 40 mm. The other steps are the same as those of specific embodiments 1 to 7.

[0041] Specific embodiment 9: This embodiment differs from specific embodiments 1 to 8 in that the time for heating the composite powder green body from room temperature to 700° C. in step 3 is 6 to 8 minutes. The other steps are the same as those of specific embodiments 1 to 8.

[0042] Specific embodiment 10: The difference between this embodiment and specific embodiments 1 to 9 is that the reaction hot pressing sintering environment described in step 3 is a vacuum environment, an argon environment or a nitrogen environment; the vacuum degree of the vacuum environment is 10 -1 ~10 Pa. The other steps are the same as those in the first to ninth embodiments.

[0043] The following examples are used to verify the beneficial effects of the present invention:

[0044] The present invention provides a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic and a preparation method thereof. The reaction formula of the component design is shown in Equation 1. When the B4C / C molar ratio is a constant value, the proportion of the ceramic phase generated by the reaction no longer changes. The Hf / Ta molar ratio can accurately regulate the ultra-high temperature boride and carbide phase composition generated by the reaction. The (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic sintered samples prepared with different Hf / Ta molar ratios in Examples 1 to 3 and Comparative Examples 1 to 2 are named H10T0B1C1-S, H9T1B1C1-S, H7T3B1C1-S, H5T5B1C1-S and H0T10B1C1-S, respectively.

[0045] 4MSi2+B4C+9C=2MB2+2MC+8SiC(M=Hf and Ta) (1)

[0046] Example 1: A method for preparing a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic (H9T1B1C1-S), which is specifically completed by the following steps:

[0047] 1. Preparation of composite powder with fine grains and uniform distribution by low-energy ball milling:

[0048] First, ZrO2 balls are added to a ball mill, and then HfSi2, TaSi2, B4C, C and anhydrous ethanol are poured into the ball mill, and then the ball mill is sealed; the sealed ball mill is fixed on a planetary ball mill, and argon is filled into the ball mill while being evacuated, so that the ball mill is subjected to low-energy ball milling for a period of time in an argon environment to obtain a ball-milled composite powder; the ball-milled composite powder is placed in a vacuum drying oven and dried at room temperature for 24 hours, and then sieved through a 140-mesh sieve in a glove box filled with argon to obtain a composite powder with refined grains and uniform distribution, and then stored in a vacuum;

[0049] The molar ratio of the total amount of HfSi2 and TaSi2 to B4C and C in step 1 is 4:1.05:9.45;

[0050] The molar ratio of HfSi2 to TaSi2 described in step 1 is 0.9:0.1;

[0051] The B4C in step 1 is in excess of 5% compared to C by the reaction formula calculation value;

[0052] The ZrO2 balls described in step 1 are large balls and small balls, the diameter of the large balls is 10 mm, the diameter of the small balls is 6 mm, and the mass ratio of the large balls to the small balls is 3:1;

[0053] The total mass ratio of the ZrO2 balls described in step 1 to the total mass of HfSi2, TaSi2, B4C and C is 10:1;

[0054] The anhydrous ethanol described in step 1 accounts for 39.47% of the total mass of HfSi2, TaSi2, B4C and C;

[0055] The rotation speed of the low-energy ball mill described in step 1 is 350 r / min, and the time of the low-energy ball mill is 24 h;

[0056] 2. Manually pre-pressing the composite powder stored in vacuum by a hydraulic press to obtain a composite powder green body;

[0057] In step 2, a graphite paper coated with a release agent is placed closely on the wall of a graphite mold, and the composite powder stored in a vacuum is poured into a circular graphite mold with a diameter of 30 mm, and then the graphite gasket and the composite powder, and the composite powder and the pressure head are separated by graphite paper coated with a release agent; the graphite mold containing the composite powder is placed under a hydraulic press and manually pressed at 3 MPa to obtain a composite powder green embryo; the release agent is talcum powder;

[0058] 3. First, the composite powder embryo is heated from room temperature to 700°C for 7 minutes; then the axial pressure of the pressure head is increased to 40MPa at a pressure increase rate of 1.1T / min, and the temperature of the composite powder embryo is increased from 700°C to 1700°C at a heating rate of 78.5°C / min during the pressure increase process. Then, it is hot-pressed and sintered for 15 minutes at 1700°C and an axial pressure of 40MPa. The sintering temperature, pressurization process and displacement curve of the powder embryo are monitored in real time by an infrared thermometer and recorded and archived at an interval of 1s. Multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics are obtained; wherein the molar ratio of Hf / Ta is 9:1, and the molar ratio of (Hf, Ta) B2 to (Hf, Ta) C is 1:1, recorded as H9T1B1C1-S;

[0059] The hot pressing sintering environment described in step 3 is a vacuum environment with a vacuum degree of 5Pa.

[0060] Example 2: The difference between this example and Example 1 is that the molar ratio of HfSi2 and TaSi2 in step 1 is 0.7:0.3; the obtained (Hf, Ta)B2-(Hf, Ta)C-SiC ceramic has a Hf / Ta molar ratio of 7:3, and a molar ratio of (Hf, Ta)B2 to (Hf, Ta)C of 1:1, which is recorded as H7T3B1C1-S. The other steps and parameters are the same as those in Example 1.

[0061] Example 3: The difference between this example and Example 1 is that the molar ratio of HfSi2 and TaSi2 in step 1 is 0.5:0.5; the molar ratio of Hf / Ta in the obtained (Hf, Ta)B2-(Hf, Ta)C-SiC ceramic is 5:5, and the molar ratio of (Hf, Ta)B2 to (Hf, Ta)C is 1:1, which is recorded as H5T5B1C1-S. The other steps and parameters are the same as those in Example 1.

[0062] Comparative Example 1: The difference between this embodiment and embodiment 1 is that the molar ratio of HfSi2 and TaSi2 in step 1 is 1:0; the molar ratio of Hf / Ta in the obtained (Hf, Ta)B2-(Hf, Ta)C-SiC ceramic is 10:0, and the molar ratio of (Hf, Ta)B2 to (Hf, Ta)C is 1:1, which is recorded as H10T0B1C1-S. The other steps and parameters are the same as those in embodiment 1.

[0063] Comparative Example 2: The difference between this embodiment and embodiment 1 is that the molar ratio of HfSi2 and TaSi2 in step 1 is 0:1; the molar ratio of Hf / Ta in the obtained (Hf, Ta)B2-(Hf, Ta)C-SiC ceramic is 0:10, and the molar ratio of (Hf, Ta)B2 to (Hf, Ta)C is 0:10, which is recorded as H0T10B1C1-S. The other steps and parameters are the same as those in embodiment 1.

[0064] The XRD patterns of the composite powders prepared with different molar ratios of HfSi2 and TaSi2 in Examples 1 to 3 and Comparative Examples 1 to 2 are as follows: Figure 2 As shown;

[0065] from Figure 2 It can be seen that the purity of the mixed powder is very high and there are almost no impurities. As the Hf / Ta molar ratio increases, the peak of the Ta-rich phase gradually increases, while the peak of the Hf-rich phase gradually decreases until it disappears.

[0066] In step 3 of Example 2, the temperature, pressure and sintering displacement data are monitored in real time to draw a sintering densification curve, such as Figure 3 As shown; Figure 3 The densification curve of H7T3B1C1-S ceramic prepared under the sintering conditions of 1700°C / 40MPa / 15min in Example 2;

[0067] The densification curve can be divided into three stages in total. The main feature of the first stage (Ⅰ) is "pressure-driven shrinkage". In this stage, the sintering displacement first tends to be flat and then drops rapidly. The reason for the decrease in displacement is that the powder particles are rearranged under pressure. The second stage (Ⅱ) can be called "temperature-driven shrinkage". In the second stage, the sintering displacement of the powder first rises slowly and then drops sharply. The volume expansion effect of heating causes a small increase in the sintering displacement. The reason for the sudden drop is that the powder particles react at high temperature, neck bonding occurs between the particles and sintering holes are formed. The volume shrinkage effect caused by the reaction causes the sintering displacement to drop rapidly. In the third stage, the sintering displacement of the powder green body decreases slightly and then gradually tends to be flat, and the volume shrinkage effect remains constant as a whole. The reason for this phenomenon is the diffusion of elements at high temperature.

[0068] Figure 4 XRD patterns of (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics with different Hf / Ta molar ratios prepared under sintering conditions of 1700°C / 40 MPa / 15 min in Examples 1-3 and Comparative Examples 1-2;

[0069] from Figure 4 It can be seen that the HfSi2 and TaSi2 diffraction peaks corresponding to the original powder do not appear in the bulk ceramics, and its main phase components are (Hf, Ta) B2, (Hf, Ta) C and SiC. The diffraction peaks of the same phase of the bulk ceramics all show a single-direction diffraction feature of high crystallinity, and no second phase appears.

[0070] Figure 5 BSE diagrams of (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics with different Hf / Ta molar ratios prepared under sintering conditions of 1700°C / 40 MPa / 15 min in Examples 1 to 3 and Comparative Examples 1 to 2, wherein (a1) and (a2) are H10T0B1C1-S, (b1) and (b2) are H9T1B1C1-S, (c1) and (c2) are H7T3B1C1-S, (d1) and (d2) are H5T5B1C1-S, and (e1) and (e2) are H0T10B1C1-S;

[0071] from Figure 5 It can be seen that the surface morphology of H10T0B1C1-S and H0T10B1C1-S ceramic samples both show loose characteristics, and a small number of sintering pores can be observed from the enlarged image ( Figure 5 (a1) and (a2), Figure 5 The surfaces of H9T1B1C1-S, H7T3B1C1-S, and H5T5B1C1-S ceramic samples are almost dense, and no sintering pores are observed in the enlarged images ( Figure 5(b1) and (b2), (c1) and (c2), (d1) and (d2)). It is also observed that with the increase of Ta element doping content, the grain size of ultra-high temperature ceramics gradually increases, and a grain interlocking structure with a large aspect ratio is formed ( Figure 5 (c1) and (c2)).

[0072] Figure 6 (a) is a TEM image of the H7T3B1C1-S ceramic prepared under the sintering conditions of 1700°C / 40MPa / 15min in Example 2; (b) is a HRTEM image of the H7T3B1C1-S ceramic prepared under the sintering conditions of 1700°C / 40MPa / 15min in Example 2;

[0073] from Figure 6 In (a), it is observed that the phases are tightly bound together, and there is no obvious gap at the grain boundary. The three regions A, B, and C are determined to be (Hf, Ta) C, (Hf, Ta) B2, and SiC, respectively, based on the electron diffraction spot results. Figure 5 In (b), it can be seen that the overall ceramic grains have good crystallinity, no amorphous phase exists, and the lattice fringes are clear. The measured interplanar spacings correspond to (Hf, Ta) C solid solution grains and SiC grains, respectively.

[0074] Figure 7 (a)-(b) are dark field images and HADDF images of H7T3B1C1-S ceramics prepared under the sintering conditions of 1700°C / 40MPa / 15min in Example 2; (c)-(h) are element distribution diagrams of H7T3B1C1-S ceramics prepared under the sintering conditions of 1700°C / 40MPa / 15min in Example 2;

[0075] Depend on Figure 7 It can be seen that Hf and Ta elements are mainly distributed in (Hf, Ta) B2 and (Hf, Ta) C phases, and Si element mainly exists in SiC phase.

[0076] Table 1 shows the mechanical properties data of (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics with different Hf / Ta molar ratios prepared in Examples 1 to 5 under the sintering conditions of 1700°C / 40MPa / 15min;

[0077] Table 2 shows the ablation performance data of (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics with different Hf / Ta molar ratios prepared in Examples 1 to 5 under the sintering conditions of 1700°C / 40MPa / 15min;

[0078] Table 1

[0079]

[0080] Table 2

[0081]

[0082] From the mechanical properties results in Table 1, it can be seen that Ta-containing solid solution ceramics exhibit better mechanical properties. The density of ceramics prepared at 1700°C for 15 minutes reaches 92%-95%. With the increase of Ta element content, the flexural strength and fracture toughness show a pattern of first increasing and then decreasing. When Hf / Ta=0.7:0.3, its flexural strength, fracture toughness and Vickers hardness are 556.61MPa and 5.13MPa·m, respectively. 1 / 2 and 25.59GPa, which are 33.96%, 58.82% and 29.64% higher than those of H10T0B1C1-S samples, respectively. From the ablation performance results in Table 2, it can be seen that the samples of H10T0B1C1-S and H9T1B1C1-S have good ablation resistance, but with the further increase of Ta element content, the ablation resistance of other samples deteriorates. When Hf / Ta=0.9:0.1, the best ablation resistance is shown, and its linear ablation rate is -0.083μm / s; the mass ablation rate is 0.095mg / s.

[0083] Figure 8 The fracture morphologies of (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics with different Hf / Ta molar ratios prepared under sintering conditions of 1700°C / 40 MPa / 15 min in Examples 1 to 3 and Comparative Examples 1 to 2, wherein (a) is H10T0B1C1-S, (b) is H9T1B1C1-S, (c) is H7T3B1C1-S, (d) is H5T5B1C1-S, and (e) is H0T10B1C1-S;

[0084] Depend on Figure 8 It can be seen that the fracture morphology of the H10T0B1C1-S sample is relatively flat, and the main fracture mode is intergranular fracture, but only a small number of small-sized grains are pulled out ( Figure 8 (a)). With the increase of Ta content, the grain size of ultra-high temperature phase increases. The fracture morphology of Ta-containing solid solution ceramics shows obvious intergranular fracture and the extraction of (Hf,Ta)B2 grains ( Figure 8 (b)-(d)). As the Ta doping increases, the oversized TaB2 grains can no longer be pulled out and obvious transgranular fracture occurs ( Figure 8 (e)).

[0085] Embodiment 4: The difference between this embodiment and embodiment 1 is that in step 1, the molar ratio of MSi2 (M = Hf and Ta, MSi2 is the sum of the molar numbers of HfSi2 and TaSi2), B4C, and C is 1:0:3.15; in step 1, the molar ratio of HfSi2 and TaSi2 is 0.9:0.1, and the obtained (Hf, Ta) C-SiC ceramic has a Hf / Ta molar ratio of 9:1, and a molar ratio of (Hf, Ta) B2 to (Hf, Ta) C is 0:1, which is recorded as H9T1B0C1-S. The other steps and parameters are the same as those in embodiment 1.

[0086] The reaction equation for the composition design is shown in Equation 2. When the HfSi2 / TaSi2 molar ratio is a constant value, the composition of the ceramic phase generated by the reaction no longer changes. Different B4C / C molar ratios can accurately control the composition ratio of the ultra-high temperature borides and carbides generated by the reaction. The sintered samples of (Hf,Ta)B2-(Hf,Ta)C-SiC ceramics with different B4C / C molar ratios were named H9T1B0C1-S, H9T1B1C2-S, H9T1B2C1-S and H9T1B1C0-S.

[0087]

[0088] Embodiment 5: The difference between this embodiment and embodiment 1 is that in step 1, the molar ratio of MSi2 (M = Hf and Ta, MSi2 is the sum of the molar numbers of HfSi2 and TaSi2), B4C, and C is 6:1.05:15.75; in step 1, the molar ratio of HfSi2 and TaSi2 is 0.9:0.1, and the obtained (Hf, Ta)B2-(Hf, Ta)C-SiC ceramic has a Hf / Ta molar ratio of 9:1, and a molar ratio of (Hf, Ta)B2 to (Hf, Ta)C is 1:2, which is recorded as H9T1B1C2-S. The other steps and parameters are the same as those in embodiment 1.

[0089] Embodiment 6: The difference between this embodiment and embodiment 1 is that in step 1, the molar ratio of MSi2 (M = Hf and Ta, MSi2 is the sum of the molar numbers of HfSi2 and TaSi2), B4C, and C is 3:1.05:6.3; in step 1, the molar ratio of HfSi2 and TaSi2 is 0.9:0.1, and the obtained (Hf, Ta)B2-(Hf, Ta)C-SiC ceramic has a Hf / Ta molar ratio of 9:1, and a molar ratio of (Hf, Ta)B2 to (Hf, Ta)C is 2:1, which is recorded as H9T1B2C1-S. The other steps and parameters are the same as those in embodiment 1.

[0090] Embodiment 7: The difference between this embodiment and embodiment 1 is that in step 1, the molar ratio of MSi2 (M = Hf and Ta, MSi2 is the sum of the molar numbers of HfSi2 and TaSi2), B4C, and C is 2:1.05:3.15; in step 1, the molar ratio of HfSi2 and TaSi2 is 0.9:0.1, and the obtained (Hf, Ta) B2-SiC ceramic has a Hf / Ta molar ratio of 9:1, and a molar ratio of (Hf, Ta) B2 to (Hf, Ta) C is 2:0, which is recorded as H9T1B1C0-S. The other steps and parameters are the same as those in embodiment 1.

[0091] Fig. 9 XRD patterns of (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics with different B4C / C molar ratios prepared under sintering conditions of 1700°C / 40MPa / 15min for Examples 4 to 7;

[0092] from Fig. 9 It can be seen that its main phase components are (Hf, Ta) B2, (Hf, Ta) C and SiC. The diffraction peaks of the same phase of bulk ceramics all show a single-direction diffraction feature of high crystallinity, and no second phase appears.

[0093] Table 3 shows the bending properties and hardness data of (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics with different B4C / C molar ratios prepared in Examples 4 to 7 under the sintering conditions of 1700°C / 40 MPa / 15 min.

[0094] Table 3

[0095]

[0096] It can be seen from Table 3 that the bending strength and hardness of H9T1B1C0-S ceramics are higher than those of H9T1B0C1-S ceramics, which can be attributed to the growth of rod-shaped crystals of solid solution boride, which consumes a lot of energy in the extraction of high aspect ratio grains during the fracture process. The same pattern is also shown in H9T1B2C1-S ceramics and H9T1B1C2-S ceramics. The bending strength and hardness of H9T1B1C2-S ceramics are higher than those of H9T1B0C1-S ceramics because the introduction of solid solution boride has a solid solution strengthening effect, which improves the mechanical properties and hardness. The increase in the performance of H9T1B2C1-S ceramics compared to H9T1B1C0-S ceramics is due to the doping of solid solution carbides.

Claims

1. A method for preparing multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramics, characterized in that: The preparation method is specifically completed according to the following steps:

1. Preparation of composite powder with fine grains and uniform distribution by low-energy ball milling: First, ZrO2 balls are added to a ball mill, and then HfSi2, TaSi2, B4C, C and anhydrous ethanol are poured into the ball mill, and then the ball mill is sealed; the sealed ball mill is fixed on a planetary ball mill, and argon is filled into the ball mill while being evacuated, so that the ball mill is subjected to low-energy ball milling for a period of time in an argon environment to obtain a ball-milled composite powder; the ball-milled composite powder is placed in a vacuum drying oven and dried at room temperature, and then sieved through a 120-200 mesh sieve in a glove box filled with argon to obtain a composite powder with refined grains and uniform distribution, and then stored in a vacuum; The molar ratio of the total amount of HfSi2 and TaSi2 to B4C and C in step 1 is (1-6):(1-1.2):(3.6-18); The molar ratio of HfSi2 to TaSi2 described in step 1 is (0.01-0.99):(0.01-0.99); 2. Manually pre-pressing the composite powder stored in vacuum by a hydraulic press to obtain a composite powder green body; 3. First, the temperature of the composite powder green body is raised from room temperature to 700 ℃; then, the axial pressure of the pressure head is increased to 30MPa~60MPa at a pressure increasing rate of 1t / min~1.3t / min. During the pressure increasing process, the temperature of the composite powder green body is raised from 700℃ to 1600℃~1800℃ at a heating rate of 76.3℃ / min~85.5℃ / min. Then, hot pressing and sintering is performed at 1600℃~1800℃ and an axial pressure of 30MPa~60MPa for 5min~15min to obtain multi-element ultra-high temperature (Hf,Ta)B2-(Hf,Ta)C-SiC ceramics.

2. The method for preparing a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic according to claim 1, characterized in that: The ZrO2 balls described in step 1 are large balls and small balls, the diameter of the large balls is 9.95mm~10.05mm, the diameter of the small balls is 5.95mm~6.05mm, and the mass ratio of the large balls to the small balls is (3~5):(1~2).

3. The method for preparing a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic according to claim 1, characterized in that: The total mass ratio of the ZrO2 balls described in step 1 to the total mass ratio of HfSi2, TaSi2, B4C and C is (6-15):

1.

4. The method for preparing a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic according to claim 1, characterized in that: The anhydrous ethanol described in step 1 accounts for 31.6% to 47.4% of the total mass of the HfSi2, TaSi2, B4C and C powders.

5. The method for preparing a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic according to claim 1, characterized in that: The rotation speed of the low-energy ball mill described in step 1 is 250r / min~350r / min, and the time of the low-energy ball milling is 24h~48h; the time of room temperature drying described in step 1 is 24h~36h.

6. The method for preparing a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic according to claim 1, characterized in that: In step 2, the graphite paper coated with a release agent is pressed against the wall of the graphite mold, the composite powder stored in vacuum is poured into the circular graphite mold, and then the graphite gasket and the composite powder, and the composite powder and the pressure head are separated by graphite paper coated with a release agent; the graphite mold containing the composite powder is placed under a hydraulic press and manually pressurized to 3MPa~5MPa to obtain a composite powder green body.

7. The method for preparing a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic according to claim 6, characterized in that: The release agent is one or more of talcum powder, pottery clay and white clay; the diameter of the graphite mold is 30mm-40mm.

8. The method for preparing a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic according to claim 1, characterized in that: In step 3, the time for heating the composite powder green body from room temperature to 700° C. is 6 min to 8 min.

9. The method for preparing a multi-element ultra-high temperature (Hf, Ta) B2-(Hf, Ta) C-SiC ceramic according to claim 1, characterized in that: The hot pressing sintering environment in step 3 is a vacuum environment, an argon environment or a nitrogen environment; the vacuum degree of the vacuum environment is 10 -1 ~10Pa.

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