High-strength wear-resistant high-entropy boride-based metal ceramic and preparation method thereof

Abstract

The invention provides a high-strength wear-resistant high-entropy boride-based metal ceramic and a preparation method thereof.The high-entropy boride metal ceramic comprises a high-entropy boride (Mo, W) 2 (Fe, Co and Ni) B2 ceramic phase and a (Fe, Co and Ni) solid solution binding phase, the two phases have good wettability, the (Mo, W) 2 (Fe, Co and Ni) B2 ceramic phase is dispersed in the continuous binding phase, and the two-phase interface has good consistency and matching degree; the two phases are firmly combined, the high-entropy boride ceramic phase has high hardness, wear resistance, high temperature resistance, oxidation resistance and chemical stability, the metal solid solution binding phase has good plasticity and toughness, and the prepared metal ceramic has excellent comprehensive performance.

Description

Technical Field

[0001] The present invention relates to the technical field of cermets, and specifically to a high-strength wear-resistant high-entropy boride-based cermet and a preparation method thereof. Background Art

[0002] Cermets are made of metal and ceramic raw materials, with the ceramic phase content being greater than 50%. Cermets not only retain the characteristics of ceramics such as high strength, high hardness, wear resistance, high temperature resistance, oxidation resistance, and chemical stability, but also possess the good toughness and flexural strength of metals. Therefore, cermets have a wide range of applications in fields such as aerospace, vehicle manufacturing, electronics industry, and tool manufacturing. The most widely used cermets include oxide-based, carbide-based, nitride-based, and boride-based cermets. Among them, boride-based cermets are considered to be one of the most promising cermets due to their high hardness, melting point, excellent electrical conductivity, and corrosion resistance.

[0003] Typical ternary boride-based cermets include Mo2FeB2-based, Mo2NiB2-based, and WCoB-based cermets. The three types of ternary boride-based cermets have their respective characteristics. For example, Mo2FeB2 cermet has excellent wear resistance, Mo2NiB2 cermet has excellent corrosion resistance, and WCoB cermet has excellent high-temperature performance. The inventor disclosed a high-entropy boride ceramic (W a Mo b )2(Fe x Co y Ni z )B2 in the previous Chinese patent document with the application number 202311635789.3. By means of high-entropy design, the components and structures of three typical ternary boride ceramics are integrated into one to form a high-entropy alloy, causing lattice distortion, thereby producing a solid solution strengthening effect and further improving the hardness of the boride ceramic. As a new type of ultra-high-temperature ceramic material, it has potential broad application prospects in fields such as aerospace and precision machining. However, there is no report on cermet materials based on this high-entropy boride ceramic. Summary of the Invention

[0004] The purpose of the present invention is to provide a high-entropy boride-based cermet and a preparation method thereof, with a simple production process and capable of preparing a high-strength wear-resistant cermet with excellent comprehensive properties.

[0005] In a first aspect, the present invention provides a high-entropy boride-based cermet, which comprises a (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and a (Fe,Co,Ni) solid-solution binder phase. Based on the mass of the high-entropy boride-based cermet, the content of the (Fe,Co,Ni) solid-solution binder phase is 17 wt.% to 23 wt.%.

[0006] Optionally, in the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase, Mo and W are solid-solved, and the atomic ratio of Mo to W is (0.4 to 0.6):(0.4 to 0.6), Fe, Co, and Ni are solid-solved, and the atomic ratio of Fe, Co, and Ni is (0.25 to 0.32):(0.33 to 0.4):(0.33 to 0.4); in the (Fe,Co,Ni) solid-solution binder phase, the atomic ratio of Fe, Co, and Ni is (0.25 to 0.32):(0.33 to 0.4):(0.33 to 0.4).

[0007] Optionally, the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase has an orthorhombic structure, and the (Fe,Co,Ni) solid-solution binder phase has a cubic structure.

[0008] Optionally, based on the mass of the high-entropy boride-based cermet, the content of the (Fe,Co,Ni) solid-solution binder phase is 18 wt.% to 21.5 wt.%.

[0009] Optionally, the grain size of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is 2 to 12 μm.

[0010] Optionally, [00-1] of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and [0-11] of the (Fe,Co,Ni) solid-solution binder phase show orientation consistency at the two-phase interface.

[0011] Optionally, the (310) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is parallel to the (311) crystal plane of the (Fe,Co,Ni) solid-solution binder phase, the (400) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is parallel to the (200) crystal plane of the (Fe,Co,Ni) solid-solution binder phase, and the (220) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is parallel to the (111) crystal plane of the (Fe,Co,Ni) solid-solution binder phase.

[0012] Optionally, the (310) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and the (311) crystal plane of the (Fe,Co,Ni) solid-solution bonding phase form a coherent interface, and the mismatch degree between the two phases is less than 50%; the (400) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and the (200) crystal plane of the (Fe,Co,Ni) solid-solution bonding phase form a coherent interface, and the mismatch degree between the two phases is less than 2.5%.

[0013] In a second aspect, the present invention provides a method for preparing the foregoing high-entropy boride-based cermet. The high-entropy boride-based cermet is prepared by mixing raw material powders composed of W powder, Mo powder, Fe powder, Co powder, Ni powder and B powder evenly, then pressing and sintering.

[0014] Optionally, the sintering temperature is 1275-1325 °C, and the sintering holding time is 0-80 min.

[0015] In summary, the present invention has at least one of the following beneficial effects:

[0016] 1. The high-entropy boride-based cermet provided by the present invention includes a (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and a (Fe,Co,Ni) solid-solution bonding phase. By selecting the specific alloy element composition and content of the bonding phase, a (Fe,Co,Ni) solid-solution bonding phase is obtained, which has good wettability to the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase, so as to form a continuous bonding phase in the ceramic matrix. During the sintering process, the bonding phase forms a liquid phase, which promotes the transformation of the intermediate phase into the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase, and at the same time promotes the dissolution and reprecipitation of the grains of the high-entropy boride ceramic phase, realizes the densification of the cermet and regulates the uniform growth of the grains of the high-entropy boride ceramic phase. The high-entropy boride ceramic phase has high hardness, wear resistance, high temperature resistance, oxidation resistance and chemical stability, and the metal solid-solution bonding phase has good plasticity and toughness. A high-strength wear-resistant cermet with excellent comprehensive performance is prepared. In subsequent embodiments, the transverse rupture strength of the cermet can reach more than 950 MPa, and the hardness can reach more than 540 MPa. In a preferred embodiment, the transverse rupture strength of the cermet can reach more than 1320 MPa, and the hardness can reach more than 840 MPa. More preferably, the transverse rupture strength of the cermet can reach more than 1530 MPa, and the hardness can reach more than 900 MPa.

[0017] 2. The high-entropy boride-based cermet provided by the present invention has a certain solubility between the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and the (Fe,Co,Ni) solid-solution binder phase, which helps the firm combination between the ceramic phase and the binder phase. At the same time, the two-phase interface has good orientation consistency along the [0-11] of the binder phase and the [00-1] of the ceramic phase. Further, the (311) crystal plane of the binder phase and the (310) crystal plane of the ceramic phase, the (200) crystal plane of the binder phase and the (400) crystal plane of the ceramic phase, and the (111) crystal plane of the binder phase and the (220) crystal plane of the ceramic phase all show good parallel relationships, which is beneficial to improving the stability and firm combination of the two-phase interface, thereby further improving the comprehensive mechanical properties of the cermet.

[0018] 3. The preparation method of the high-entropy boride-based cermet provided by the present invention is to mix the raw material powders composed of W powder, Mo powder, Fe powder, Co powder, Ni powder and B powder evenly and then prepare the high-entropy boride-based cermet by compacting and sintering. It is prepared by the one-step reaction sintering of elemental raw material powders. The process is simple and easy to control, avoiding the complicated processes of separately prefabricating the hard phase and the binder phase. Moreover, the prepared cermet has a stable structure, which is beneficial to improving the production efficiency and the product yield. Description of the Drawings

[0019] Figure 1 . XRD pattern of the sample sintered between 1275 °C and 1325 °C in Example 1.

[0020] Figure 2 . Microstructures of the samples sintered at different temperatures in Example 1: (a) 1275 °C, (b) 1290 °C, (c) 1300 °C, (d) 1310 °C, (e) 1325 °C.

[0021] Figure 3 . EDS analysis of the sample sintered at 1300 °C in Example 1: (a) ceramic phase, (b) binder phase.

[0022] Figure 4 . HADDF and EDS mapping images of the sample sintered at 1300 °C in Example 1.

[0023] Figure 5 . HRTEM images and FFT patterns of the sample sintered at 1300 °C in Example 1: (a) FFT of the binder phase, (b) FFT of the ceramic phase, (c) FFT of the interface.

[0024] Figure 6 . XRD pattern of the samples with different sintering holding times at 1310 °C in Example 2.

[0025] Figure 7Microstructures of samples with different sintering holding times at 1310 °C in Example 2: (a) 0 minute, (b) 20 minutes, (c) 40 minutes, (d) 60 minutes, (e) 80 minutes.

[0026] Figure 8 XRD patterns of cermets with different binder phase contents in Example 3.

[0027] Figure 9 Microstructures of samples with different binder phase contents in Example 3: (a) 17 wt.%, (b) 18.5 wt.%; (c) 20 wt.%, (d) 21.5 wt.%, (e) 23 wt.%. Detailed implementation manners

[0028] The present invention provides a high-strength wear-resistant high-entropy boride-based cermet and a preparation method thereof. To make the objectives, technical solutions and effects of the present invention clearer and more definite, the present invention is further described in detail below. It should be understood that the specific embodiments described herein are only used to explain the present invention and are not used to limit the present invention.

[0029] Typical ternary boride-based cermets, such as Mo2FeB2-Fe, Mo2NiB2-Ni, and WCoB-Co, combine the excellent properties of borides and metals and show great development potential in the fields of wear resistance, corrosion resistance, and high temperature resistance, and have become a research hotspot in the field of cemented carbide materials. However, with the rapid development of modern industry, these cermets gradually show deficiencies in practical applications. For example, the Mo2FeB2-Fe cermet has insufficient corrosion resistance, the Mo2NiB2-Ni cermet has poor wear resistance, and the overall mechanical properties of the WCoB-Co cermet are insufficient. The inventor disclosed a novel high-entropy ternary boride ceramic (W a Mo b )2(Fe x Co y Ni z)B2 combines the components and structures of three typical ternary boride-based ceramics. W and Mo are solid-soluted, and Fe, Co, and Ni are solid-soluted to form a high-entropy alloy, causing lattice distortion and thus producing a solid-solution strengthening effect. On this basis, through further research, the inventor prepared a new type of high-entropy boride-based cermet. The cermet includes a (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and a (Fe,Co,Ni) solid-solution bonding phase. By selecting the specific alloy element composition and content of the bonding phase, a (Fe,Co,Ni) solid-solution bonding phase with good wettability to the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is obtained, thereby forming a continuous bonding phase in the ceramic matrix, forming a liquid phase during sintering, promoting the transformation of the intermediate phase to the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase, and at the same time promoting the dissolution-reprecipitation of the grains of the high-entropy boride ceramic phase, realizing the densification of the cermet and the controllable and uniform growth of the grains of the high-entropy boride ceramic phase. In addition, the high-entropy boride ceramic phase and the solid-solution bonding phase show good orientation consistency at the two-phase interface, which is beneficial to improving the stability of the combination of the two phases, thereby further improving the comprehensive mechanical properties of the cermet. The present invention is obtained on the basis of this research.

[0030] In some embodiments of the present invention, the high-entropy boride-based cermet includes a (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and a (Fe,Co,Ni) solid-solution bonding phase. Based on the mass of the high-entropy boride-based cermet, the content of the (Fe,Co,Ni) solid-solution bonding phase is 17 wt.% to 23 wt.%. Preferably, the content of the (Fe,Co,Ni) solid-solution bonding phase is 18 wt.% to 21.5 wt.%. More preferably, the content of the (Fe,Co,Ni) solid-solution bonding phase is 18.5 wt.% to 21.5 wt.%.

[0031] In some embodiments of the present invention, in the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase, Mo and W are solid-soluted, and the atomic ratio of Mo to W is (0.4 to 0.6):(0.4 to 0.6), Fe, Co, and Ni are solid-soluted, and the atomic ratio of Fe, Co, and Ni is (0.25 to 0.32):(0.33 to 0.4):(0.33 to 0.4); in the (Fe,Co,Ni) solid-solution bonding phase, the atomic ratio of Fe, Co, and Ni is (0.25 to 0.32):(0.33 to 0.4):(0.33 to 0.4).

[0032] In some embodiments of the present invention, the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase has an orthorhombic structure, and the (Fe,Co,Ni) solid-solution binder phase has a cubic structure.

[0033] In some embodiments of the present invention, the grain size of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is 2 - 12 μm, preferably 2 - 10 μm, and more preferably 2 - 8 μm.

[0034] In some embodiments of the present invention, [00-1] of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and [0-11] of the (Fe,Co,Ni) solid-solution binder phase show orientation consistency at the two-phase interface. The (310) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is parallel to the (311) crystal plane of the (Fe,Co,Ni) solid-solution binder phase, the (400) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is parallel to the (200) crystal plane of the (Fe,Co,Ni) solid-solution binder phase, and the (220) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is parallel to the (111) crystal plane of the (Fe,Co,Ni) solid-solution binder phase. Optionally, the (310) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and the (311) crystal plane of the (Fe,Co,Ni) solid-solution binder phase form a coherent interface, and the mismatch degree between the two phases is less than 50%; the (400) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and the (200) crystal plane of the (Fe,Co,Ni) solid-solution binder phase form a coherent interface, and the mismatch degree between the two phases is less than 2.5%.

[0035] In some embodiments of the present invention, the high-entropy boride-based cermet is prepared by mixing the raw material powders composed of W powder, Mo powder, Fe powder, Co powder, Ni powder and B powder evenly and then pressing and sintering. Optionally, the sintering temperature is 1275 - 1325 °C, preferably 1290 - 1325 °C, and more preferably 1290 - 1315 °C; optionally, the sintering holding time is 0 - 80 min, preferably 20 - 80 min, and more preferably 30 - 70 min.

[0036] In some embodiments of the present invention, the high-entropy boride-based cermet has a transverse rupture strength of 950 to 1650 MPa and a hardness of 545 to 970 MPa; preferably, the transverse rupture strength is 1190 to 1650 MPa and the hardness is 840 to 970 MPa; more preferably, the transverse rupture strength is 1320 to 1650 MPa and the hardness is 840 to 970 MPa; more preferably, the transverse rupture strength is 1500 to 1650 MPa and the hardness is 840 to 970 MPa; more preferably, the transverse rupture strength is 1590 to 1650 MPa and the hardness is 890 to 970 MPa.

[0037] The following further describes the present invention in detail with specific examples and comparative examples. In the specific examples of the present invention, the raw materials are commercially available unless otherwise specified. The particle size of each raw material powder is below 9 μm and the purity is above 99%. The specific particle size and specifications are as follows: W powder: average particle size ~2 μm, purity ≥99.9%; Mo powder: average particle size ~2 μm, purity ≥99.9%; Fe powder: average particle size 3 - 5 μm, purity ≥99.5%; Co powder: average particle size 1 - 3 μm, purity ≥99.5%; Ni powder: average particle size 1 - 3 μm, purity ≥99.5%; B powder: average particle size ~1 μm, purity ≥99.3%.

[0038] In the following examples, the chemical formulas of each phase represent the components and the atomic ratios of each component. For example, the chemical formula of the ceramic phase is (W 0.5 Mo 0.5 )2(Fe 0.3 Co 0.35 Ni 0.35 )B2, indicating the atomic ratios of each component as follows: W:Mo:Fe:Co:Ni:B = (0.5×2):(0.5×2):0.3:0.35:0.35:2 = 1:1:0.3:0.35:0.35:2; the chemical formula of the binder phase is (Fe 0.3 Co 0.35 Ni 0.35 ), indicating the atomic ratios of each component as follows: Fe:Co:Ni = 0.3:0.35:0.35. The 20 in the nominal composition of the cermet (W 0.5 Mo 0.5 )2(Fe 0.3 Co 0.35 Ni 0.35 )B2 - 20(Fe 0.3 Co 0.35 Ni 0.35 ) means that the (Fe 0.3 Co 0.35 Ni 0.35 ) binder phase is 20 wt% of the total mass of the cermet (ceramic phase + binder phase).

[0039] Example 1

[0040] The nominal composition of the cermet prepared in this example is (W 0.5 Mo 0.5 )2(Fe 0.3 Co 0.35 Ni 0.35 )B2-20(Fe 0.3 Co 0.35 Ni 0.35 ). The specific preparation process steps are as follows:

[0041] 1). Weigh the raw material powders according to the above nominal composition ratio. The raw material powders are W powder, Mo powder, Fe powder, Co powder, Ni powder and B powder;

[0042] 2). Ball-mill and mix the above raw material powders. Wet milling is used for ball milling. The ball milling medium is cemented carbide balls, the ball-to-material ratio is 8:1, the ball milling solvent is anhydrous ethanol, the addition ratio of anhydrous ethanol to the raw material powders is 1.2 ml:1 g, the ball milling speed is 100 rpm, and the stirring time is 72 h. After the ball milling is completed, take out the mixed slurry and dry it in vacuum at 50 °C to obtain the mixed powder;

[0043] 3). Sinter the mixed powder. Press the mixed powder into a mold under a pressure of 150 MPa, and then carry out vacuum sintering. The sintering vacuum degree is ≤10 Pa. The sintering temperature is sintered at 5 temperatures in Table 1 respectively, and the sintering holding time is 60 min; after the sintering is completed, cool it with the furnace and take it out to prepare samples A-1, A-2, A-3, A-4, A-5.

[0044] Table 1. Different sintering temperatures corresponding to 5 samples

[0045]

[0046]

[0047] The phase composition of the cermet prepared at different sintering temperatures in Example 1 was tested by X-ray diffraction (XRD, SHIMADZU XRD-7000) using Cu-Kα radiation. The test results are as Figure 1 shown. From Figure 1It can be seen that the phase structures of the samples sintered at 1275℃ to 1325℃ in Example 1 are (Mo,W)2(Fe,Co,Ni)B2, (Fe,Co,Ni) and a small amount of (Fe,Co,Ni)7(Mo,W)6 phases. This shows that within this temperature range, the reaction in the system has basically occurred completely. When the sintering temperature is increased from 1275℃ to 1325℃, the relative intensity of the diffraction peak of the (Fe,Co,Ni)7(Mo,W)6 phase decreases from 4.6% to 2.7%. This phenomenon may be attributed to the increase in liquid phase content and the enhancement of atomic diffusion capacity caused by the increase in sintering temperature, and more Mo and W atoms are dissolved in the bonding phase, resulting in a reduction in the (Fe,Co,Ni)7(Mo,W)6 phase.

[0048] The microstructure and element distribution of the sintered sample of Example 1 were tested by scanning electron microscopy (SEM, JEOL JSM-6510LV) and energy dispersive spectroscopy (EDS). Figure 2 and Figure 3 As shown. Figure 2 It can be seen that when sintered at 1275°C, more pores are observed in the sample. This is because the sintering temperature at this time is relatively low and the liquid phase content is low. As the sintering temperature increases, the liquid phase continues to form and fill the pores. At 1290°C, the number of pores decreases significantly. When the temperature rises to 1300°C, the sample achieves initial densification. In addition, at 1300°C, some ceramic phase grains are observed to agglomerate. This is mainly because at this temperature, the liquid phase is not fully diffused, and the rearrangement and dissolution and re-precipitation effects of the ceramic phase grains are not fully carried out, resulting in a certain degree of inhomogeneity in the sample microstructure. When the temperature rises to 1310°C, the agglomerated particles are significantly reduced, and the sample presents a more uniform microstructure. As the sintering temperature continues to rise to 1325°C, it is observed that some ceramic phase grains coarsen and begin to merge to form larger grains. Combined with XRD (such as Figure 1 ) and EDS analysis of samples sintered at 1300℃ (as shown Figure 3 shown), Figure 2 The gray-white phase was identified as the (Mo, W)2(Fe, Co, Ni)B2 ceramic phase, and the gray-black phase was identified as the (Fe, Co, Ni) solid solution bonding phase. Combined with the EDS analysis results, it can be seen that there is a certain solubility between the ceramic phase and the bonding phase, which helps to firmly bond the two phases. Figure 2It can also be observed that the (Mo,W)2(Fe,Co,Ni)B2 ceramic phase appears as granular or lath-shaped, distributed in the continuous phase (Fe,Co,Ni) solid solution. The grain size of the ceramic phase is 2-8 μm, and the (Mo,W)2(Fe,Co,Ni)B2 ceramic phase is closely bonded to the (Fe,Co,Ni) solid solution. When the temperature reaches 1325 °C, there is an obvious coarsening phenomenon in the ceramic phase, and the maximum grain size is greater than 8 μm, but still remains within 10 μm.

[0049] A three-point bending instrument was used to evaluate and measure the transverse rupture strength (TRS, MST CMT5504) of the samples, and the reference test standard was GB / T 3851-2015 "Determination Method for Transverse Rupture Strength of Cemented Carbides". A Vickers hardness tester (HV, HV-30Z) was used to measure the hardness of the samples, and the reference test standard was GB / T 7997-2014 "Test Method for Vickers Hardness of Cemented Carbides". The test results are shown in Table 2.

[0050] Table 2. Mechanical properties of 5 samples

[0051] Sample A-1 A-2 A-3 A-4 A-5 Sintering temperature 1275℃ 1290℃ 1300℃ 1310℃ 1325℃ Hardness / MPa 547.6 910.4 938.9 902.4 877.3 TRS / MPa 959.4 1191.3 1326.5 1639.2 1507.9

[0052] Table 2 shows the trends of the hardness and transverse rupture strength (TRS) of the cermet samples sintered in vacuum for 60 min in Example 1 with the change of sintering temperature. It can be seen from Table 2 that with the increase of sintering temperature, the hardness and transverse rupture strength of the cermet first increase and then show a downward trend. The main reason for the analysis is that when the sintering temperature is 1275 °C, the sintering behavior and liquid-phase sintering of the system are insufficient, and the lower relative density results in poor transverse rupture strength and hardness of the cermet; with the increase of temperature, the liquid-phase sintering process continues. At 1290 °C, the relative density of the specimen increases, and at the same time, the transverse rupture strength and hardness increase significantly; when the temperature rises to 1300 °C, the cermet reaches a relatively dense state, and the transverse rupture strength and hardness increase to 1326.5 MPa and 938.9 MPa respectively. However, the uneven microstructure caused by the insufficient diffusion of the liquid phase will cause local stress concentration when the specimen is stressed, which is not conducive to further improving the transverse rupture strength of the sample; when the sintering temperature rises to 1310 °C, the liquid phase diffuses sufficiently, and the sample shows a more uniform microstructure, and the transverse rupture strength of the cermet also reaches the highest value of 1639.2 MPa at this temperature; when the temperature further rises to 1325 °C, the ceramic phase grains show a coarsening phenomenon, resulting in a decrease in the hardness and transverse rupture strength of the sample.

[0053] The microstructure of the sample sintered at 1300 °C in Example 1 was tested using a transmission electron microscope (TEM, FEI Tecnai G2 F20) and energy dispersive spectroscopy (EDS, X-Twin), and the test results are as Figure 4 and Figure 5 shown. Figure 4 Figure Figure 4 shows the high-angle annular dark field (HAADF) and corresponding EDS mapping images of the sample sintered at 1300 °C. The EDS results indicate that the distributions of Fe, Co, and Ni are consistent and concentrated in the binder phase, while the distributions of W, Mo, and B are similar and concentrated in the ceramic phase. The cermet exhibits a typical two-phase structure, with an obvious phase interface between the ceramic phase and the binder phase, and the two phases are firmly combined. High-resolution transmission electron microscopy (HRTEM) was used to characterize the interfacial features of the ceramic phase and the binder phase, and the results are as Figure 5 shown. The fast Fourier transform (FFT) patterns corresponding to the ceramic phase (white square), binder phase (orange square), and interface (red square) are as Figure 5 (a), (b), (c) shown, Figure 5 (a) shows that the binder phase exhibits a typical cubic structure, Figure 5 (b) shows that the ceramic phase exhibits a typical orthorhombic structure, Figure 5 (c) shows that the interface between the ceramic phase and the binder phase exhibits good orientation consistency, with the binder phase in the [0-11] direction and the ceramic phase in the [00-1] direction. In addition, (311) 粘结相 and (310) 陶瓷相 , (200) 粘结相 and (400) 陶瓷相 , (111) 粘结相 and (220) 陶瓷相 all exhibit good parallel relationships. At the same time, the measured lattice spacings of (200) 粘结相 and (400) 陶瓷相 are 1.776 nm and 1.733 nm, respectively, and the mismatch degree between the two phases is 2.4%, which is a typical coherent interface. In addition, the lattice spacings of (311) 粘结相 and (310) 陶瓷相 are 1.061 nm and 2.030 nm, respectively, and there is a significant mismatch degree of 47.7% between the two phases, which is also a typical coherent interface. However, since the lattice spacings between (311) 粘结相 and (310) 陶瓷相 are almost twice different, the elastic strain energy of the interface can be effectively reduced, thereby improving the interface stability.

[0054] Example 2

[0055] The difference between Example 2 and Example 1 is that, 3). In the mixed powder sintering step, the sintering temperature is 1310°C, and the sintering holding time is carried out according to the 4 holding times in Table 3; the remaining preparation steps are the same as those in Example 1, and samples A-6, A-7, A-8, and A-9 are prepared. For the convenience of analysis, sample A-4 sintered for 60 minutes in Example 1 is added for comparison.

[0056] Table 3. Sintering at 1310℃ with 5 different holding times

[0057] Sample A-6 A-7 A-8 A-4 A-9 Sintering holding time 0 min 20 min 40 min 60 min 80 min

[0058] The same test method as in Example 1 was used to test the phase structure of the sintered sample of Example 2. The XRD results are as follows: Figure 6 As shown, from Figure 6 It can be seen that the extension of the holding time did not significantly change the phase composition of the sample. The diffraction peaks of (Mo, W)2(Fe, Co, Ni)B2 phase, (Fe, Co, Ni) phase and a small amount of (Fe, Co, Ni)7(Mo, W)6 phase were detected in all samples.

[0059] The microstructure of the sintered sample of Example 2 was tested using the same test method as Example 1. The SEM results are as follows: Figure 7 As shown in the figure, when the sintering system does not have a holding stage, it can be observed that there are many pores on the surface of the sample, the bonding phase is unevenly distributed, and the ceramic phase grains have obvious agglomeration. This may be because the liquid phase does not flow sufficiently, and the filling of the pores and the rearrangement of the ceramic phase grains are not fully carried out; with the extension of the holding time, the liquid phase continuously fills the pores, promoting the dissolution and rearrangement of the ceramic phase grains. When the holding time is 20 minutes, the number and size of the pores on the sample surface are significantly reduced, and the agglomeration of the particles is improved; when the holding time is extended from 20 minutes to 40 minutes, the pores on the sample surface almost disappear, and the sample presents a relatively uniform microstructure; when the holding time is 60 minutes, it can be observed that some ceramic phase grains are coarsened, and the ceramic phase grain size is maintained at 2-8μm; when the holding time is further extended to 80 minutes, the grain coarsening phenomenon is obvious, and some grains begin to merge to form larger grains. The maximum grain size exceeds 8μm, but still remains within 10μm.

[0060] The same test method as in Example 1 was used to test the performance of the sample in Example 2. The results are shown in Table 4.

[0061] Table 4. Properties of samples at different sintering holding times at 1310℃

[0062]

[0063] Table 4 shows the trend of hardness and transverse rupture strength of metal ceramic samples sintered at 1310℃ vacuum over time. It can be seen from Table 4 that with the increase of the holding time of the sintering system, the transverse rupture strength of the metal ceramic samples first gradually increases and then decreases; while the hardness shows a monotonically decreasing trend with the increase of the holding time. When the system does not set a holding stage, the liquid phase is not fully diffused, and the presence of more pores leads to a low relative density, which deteriorates the strength of the metal ceramic. When the holding time is 20min, the density of the sample increases and the strength also increases accordingly. When the holding time increases to 40min, the pores are significantly reduced, the agglomeration of particles is significantly improved, and the more uniform microstructure further improves the transverse rupture strength of the sample. When the holding time is 60min, some grains coarsen, but at this time, the contribution of the homogenization of the microstructure to the transverse rupture strength occupies a dominant role, which further increases the strength of the sample to the maximum value. Continue to extend the holding time, the ceramic phase grains are significantly coarsened, resulting in a decrease in the transverse rupture strength of the sample. The hardness of metal ceramics is generally affected by multiple factors such as porosity, uniformity of organization, and grain size. As the holding time increases from 0 min to 80 min, the hardness of the metal ceramic sample gradually decreases, indicating that as the holding time increases, grain growth plays a major role. In summary, the high entropy boride-based metal ceramics prepared at a sintering temperature of 1310 ° C in vacuum for 40 min have more excellent comprehensive mechanical properties, and their transverse rupture strength and hardness are 1611.2 MPa and 936.0 MPa, respectively.

[0064] Example 3

[0065] The difference between Example 3 and Example 2 is that the content of the binder phase in the nominal composition of the metal ceramic of Example 3 is different from that of Example 2. The content of the binder phase is shown in Table 5. The specific preparation process steps are as follows:

[0066] 1). Weigh the raw material powders according to the above nominal composition ratio, the raw material powders are W powder, Mo powder, Fe powder, Co powder, Ni powder and B powder;

[0067] 2) The raw material powders are mixed by ball milling. The ball milling adopts wet milling. The ball milling medium is carbide balls. The ball-to-material ratio is 8:1. The ball milling solvent is anhydrous ethanol. The addition ratio of anhydrous ethanol to raw material powder is 1.2ml:1g. The ball milling speed is 100rpm. The stirring time is 72h. After the ball milling is completed, the mixed slurry is taken out and vacuum dried at 50℃ to obtain a mixed powder;

[0068] 3). Sintering of the mixed powder: The mixed powder is pressed into a shape under a pressure of 150 MPa, and then vacuum sintered. The sintering vacuum degree is ≤10 Pa, the sintering temperature is 1310 °C, and the sintering holding time is 40 min. After sintering, it is cooled in the furnace and taken out to prepare samples A-10, A-11, A-12, and A-13. For the convenience of analysis, the sintered sample A-8 with a binder phase of 20 wt% in the nominal composition of the cermet in Example 2 is added.

[0069] Table 5. Sintered samples with different binder phase contents sintered in vacuum at 1310 °C for 40 min

[0070] Sample A-10 A-11 A-8 A-12 A-13 Binder phase content 17 wt.% 18.5 wt.% 20 wt.% 21.5 wt.% 23 wt.%

[0071] The phase structure of the sintered samples in Example 3 was tested using the same test method as in Example 2. The XRD results are as Figure 8 shown. It can be seen from Figure 8 that with the increase of the binder phase content, the phase composition of the cermet did not change significantly. The main phase (Mo,W)2(Fe,Co,Ni)B2 ceramic phase and (Fe,Co,Ni) solid solution binder phase, as well as a small amount of (Fe,Co,Ni)7(Mo,W)6 phase, were detected in all samples.

[0072] The microstructure of the sintered samples in Example 3 was tested using the same test method as in Example 2. The SEM results are as Figure 9 shown. When the binder phase content is 17 wt.%, there is a certain agglomeration phenomenon in the ceramic phase, which may be because the binder phase content is low and the generated liquid phase is not sufficient to flow fully between the ceramic phase grains. When the binder phase content is 18.5 wt.%, the agglomeration phenomenon of the particles is improved, and the distribution of the ceramic phase and the binder phase is relatively uniform. With the further increase of the binder phase content, the increase of the liquid phase content promotes the liquid phase sintering process, and the dissolution and precipitation process intensifies, resulting in the gradual growth of the ceramic phase grains, and the grain size is 2 - 8 μm. Especially when the binder phase content is 23 wt.%, abnormally grown grains are observed in the sample, and the maximum grain size still remains within 12 μm.

[0073] Using the same test method as in Example 2, the performance of the samples in Example 3 was tested, and the results are shown in Table 6.

[0074] Table 6. Performance of samples with different binder phase contents sintered in vacuum at 1310 °C for 40 min

[0075] Sample A-10 A-11 A-8 A-12 A-13 Hardness / MPa 1020.1 969.7 936.0 896.1 848.7 TRS / MPa 1340.1 1596.6 1611.2 1604.1 1519.9

[0076] Table 6 describes the influence of the binder phase content on the mechanical properties of cermets. It can be seen that as the binder phase content increases, the transverse rupture strength of the cermet first increases and then shows a decreasing trend; while the hardness gradually decreases with the increase of the binder phase content. When the binder phase content is 17 wt.%, due to a certain degree of agglomeration of the ceramic phase, the transverse rupture strength of the sample is relatively low. When the binder phase content is increased to 18.5 wt.%, the sample shows a more uniform microstructure, and the transverse rupture strength of the cermet is significantly improved. When the binder phase content is increased to 20 wt.%, the transverse rupture strength of the sample reaches the highest value. When the binder phase content gradually increases from 20 wt.% to 23 wt.%, due to the increasingly obvious grain coarsening phenomenon, the strength shows a downward trend. In terms of hardness, since in cermets, the metal binder phase is a soft phase, the higher its content, the lower the overall hardness of the cermet material. Therefore, as the binder phase content increases, the overall hardness of the cermet gradually decreases.

[0077] Comparative Example 1

[0078] The nominal composition of the cermet prepared in Comparative Example 1 is Mo2FeB2-18.5Fe. The chemical formula of the Mo2FeB2 ceramic phase represents the atomic ratios of the components as follows: Mo:Fe:B = 2:1:2; 18.5 indicates that the Fe binder phase is 18.5 wt% of the total mass of the cermet (ceramic phase + binder phase). The specific preparation process steps are as follows:

[0079] 1). Weigh the raw material powders according to the above nominal composition ratio. The raw material powders are Mo powder, Fe powder and B powder;

[0080] 2). Ball-mill and mix the above raw material powders. Wet ball-milling is used, the ball-milling medium is cemented carbide balls, the ball-to-powder ratio is 8:1, the ball-milling solvent is anhydrous ethanol, the addition ratio of anhydrous ethanol to the raw material powders is 1.2 ml:1 g, the ball-milling rotation speed is 100 rpm, the stirring time is 72 h. After the ball-milling is completed, take out the mixed slurry and dry it in vacuum at 50 °C to obtain the mixed powder;

[0081] 3). Sinter the mixed powder. Press the mixed powder into a mold under a pressure of 150 MPa, and then carry out vacuum sintering. The sintering vacuum degree is ≤10 Pa, the sintering temperature is 1310 °C, and the sintering holding time is 40 min; after the sintering is completed, cool it with the furnace and take it out to prepare sample B-1.

[0082] Comparative Example 2

[0083] The chemical formula of the cermet prepared in Comparative Example 2 is Mo2NiB2-18.5Ni. The chemical formula of the Mo2NiB2 ceramic phase represents the atomic ratios of the components as follows: Mo:Ni:B = 2:1:2; 18.5 indicates that the Ni binder phase is 18.5 wt% of the total mass of the cermet (ceramic phase + binder phase). The specific preparation process steps are as follows:

[0084] 1). Weigh the raw material powders according to the above nominal composition ratio. The raw material powders are Mo powder, Ni powder, and B powder;

[0085] 2). Ball-mill and mix the above raw material powders. Wet ball-milling is used, the ball-milling medium is cemented carbide balls, the ball-to-material ratio is 8:1, the ball-milling solvent is anhydrous ethanol, the addition ratio of anhydrous ethanol to the raw material powders is 1.2 ml:1 g, the ball-milling rotation speed is 100 rpm, the stirring time is 72 h. After the ball-milling is completed, take out the mixed slurry and vacuum-dry it at 50 °C to obtain the mixed powder;

[0086] 3). Sinter the mixed powder. Press the mixed powder into a mold under a pressure of 150 MPa, and then carry out vacuum sintering. The sintering vacuum degree is ≤10 Pa, the sintering temperature is 1310 °C, and the sintering holding time is 40 min; after the sintering is completed, cool it with the furnace and take it out to prepare sample B-2.

[0087] Using the same test method as in Example 1, test the mechanical properties of the samples of Comparative Example 1 and Comparative Example 2 sintered in vacuum at 1310 °C for 40 min. The test results are shown in Table 7. For the convenience of analysis, sample A-11 with a binder phase of 18.5 wt.% in the nominal composition of the cermet in Example 3 is added.

[0088] Table 7. Properties of cermets with different nominal compositions

[0089]

[0090] The mechanical properties of different ternary boride systems are shown in Table 7. It can be seen from Table 7 that the Mo2FeB2-18.5Fe cermet has the highest hardness but a lower transverse rupture strength, while the Mo2NiB2-18.5Ni has a higher transverse rupture strength but a lower hardness. Comparatively speaking, the high-entropy boride-based cermet has more excellent comprehensive mechanical properties.

[0091] It should be understood that the application of the present invention is not limited to the above examples. For those of ordinary skill in the art, improvements or transformations can be made according to the above description. All these improvements and transformations should fall within the protection scope of the appended claims of this application.

Claims

1. A high-entropy boride-based cermet, characterized in that, The high-entropy boride-based cermet includes a (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and a (Fe,Co,Ni) solid-solution bonding phase. Based on the mass of the high-entropy boride-based cermet, the content of the (Fe,Co,Ni) solid-solution bonding phase is 17 wt.% to 23 wt.%.

2. The high-entropy boride-based cermet according to claim 1, wherein In the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase, Mo and W are solid-solved, and the atomic ratio of Mo to W is (0.4 - 0.6):(0.4 - 0.6). Fe, Co, and Ni are solid-solved, and the atomic ratio of Fe, Co, and Ni is (0.25 - 0.32):(0.33 - 0.4):(0.33 - 0.4). In the (Fe,Co,Ni) solid-solution bonding phase, the atomic ratio of Fe, Co, and Ni is (0.25 - 0.32):(0.33 - 0.4):(0.33 - 0.4).

3. The high-entropy boride-based cermet according to claim 1 or 2, characterized in that, The (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase has an orthorhombic structure, and the (Fe,Co,Ni) solid-solution bonding phase has a cubic structure.

4. The high-entropy boride-based cermet according to claim 1 or 2, characterized in that, Based on the mass of the high-entropy boride-based cermet, the content of the (Fe,Co,Ni) solid-solution bonding phase is 18 wt.% to 21.5 wt.%.

5. The high-entropy boride-based cermet according to claim 1 or 2, characterized in that, The grain size of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is 2 - 12 μm.

6. The high-entropy boride-based cermet according to claim 1 or 2, characterized in that, The [00-1] of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and the [0-11] of the (Fe,Co,Ni) solid-solution bonding phase show orientation consistency at the two-phase interface.

7. The high-entropy boride-based cermet according to claim 6, characterized in that, The (310) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is parallel to the (311) crystal plane of the (Fe,Co,Ni) solid-solution bonding phase. The (400) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is parallel to the (200) crystal plane of the (Fe,Co,Ni) solid-solution bonding phase. The (220) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase is parallel to the (111) crystal plane of the (Fe,Co,Ni) solid-solution bonding phase.

8. The high-entropy boride-based cermet according to claim 6, characterized in that, The (310) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and the (311) crystal plane of the (Fe,Co,Ni) solid-solution bonding phase form a coherent interface, and the mismatch degree between the two phases is less than 50%. The (400) crystal plane of the (Mo,W)2(Fe,Co,Ni)B2 high-entropy boride ceramic phase and the (200) crystal plane of the (Fe,Co,Ni) solid-solution bonding phase form a coherent interface, and the mismatch degree between the two phases is less than 2.5%.

9. A method for preparing the high-entropy boride-based cermet according to any one of claims 1 to 8, characterized in that, The high-entropy boride-based cermet is prepared by mixing the raw material powders composed of W powder, Mo powder, Fe powder, Co powder, Ni powder, and B powder evenly, followed by compacting and sintering.

10. The preparation method of the high-entropy boride-based cermet according to claim 9, characterized in that, The sintering temperature is 1275 to 1325 °C, and the holding time for sintering is 0 to 80 min.

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