A binder-free cemented carbide and its preparation method and application

Abstract

The present invention discloses a binderless cemented carbide, a preparation method, and applications thereof, relating to the technical field of cemented carbide. Because the binder metal content in the binderless cemented carbide provided by the present invention is less than or equal to 1 wt.%, WC grains are difficult to grow into a plate-like morphology by dissolving and precipitating in the binder metal. The inventors, by adding elements such as boron, titanium, chromium, and vanadium and controlling their dosage, achieve directionally controlled diffusion growth rates at the grain boundaries of each WC grain plane, thereby controlling the growth of some grains to form polygonal flaky morphologies. This allows the preparation of polygonal flaky grains in the absence of a binder metal and without requiring flattening of the raw material powder. This effectively improves the toughness of the binderless cemented carbide without significantly reducing its hardness and thus its wear resistance.

Description

Technical Field

[0001] The present invention relates to the technical field of cemented carbide, in particular to a binder-free cemented carbide and a preparation method and application thereof. Background Art

[0002] Traditional WC-Co (Ni, Fe)-other additive cemented carbides have been widely used in many fields such as cutting tools, molds, mining tools, and wear-resistant parts due to their good comprehensive properties such as strength, hardness, toughness and wear resistance. Traditional cemented carbides are mainly prepared by powder metallurgy, that is, WC powder, binder phase powder (mainly including Co, Ni or Fe, whose weight content is generally 6wt.% to 25wt.%), and other additive powders are mixed by ball milling, molded and sintered. During the sintering process, fine WC particles (or grains) dissolve in the binder liquid phase, and then precipitate and grow on the undissolved WC grains. Since WC is a close-packed hexagonal structure (HCP) crystal, it is easy to grow preferentially, so the WC grains in traditional cemented carbides are usually triangular prisms.

[0003] To improve the hardness and toughness of conventional cemented carbide, those skilled in the art have exploited the anisotropic characteristics of HCP-structured WC, where the hardness of the {0001} end crystal plane is much higher than that of the {1100} prismatic crystal plane. This allows the WC grain morphology to be transformed from the usual triangular prism to a plate (sheet) shape. This promotes the lateral growth of the WC grain prism plane and increases the proportion of the {0001} crystal plane, thereby obtaining a cemented carbide with both high hardness and toughness. To obtain plate (sheet)-shaped WC grains, it is usually necessary to use flattened raw material powder, such as tungsten powder, tungsten carbide powder, or tungsten-cobalt-carbon compound powder. During sintering, the growth rate along the flattening direction is greater than the growth rate in the perpendicular direction, thus forming plate (sheet)-shaped WC grains. However, the above schemes have the following shortcomings: (1) The flattening process of the raw material powder is complicated; (2) Since the plate-like grains formed during sintering also grow in the vertical direction, the thickness of the plate-like grains is often thicker (generally greater than 0.8 μm); (3) Due to the dissolution-precipitation growth mechanism, almost all plate-like grains are formed, and it is difficult to quantitatively control the content of plate-like grains.

[0004] In addition, due to the inherent characteristics of high binder phase metal content, traditional cemented carbide has good toughness and fracture toughness K 1C Generally higher than 7MPa·m 1 / 2, but its hardness (Vickers hardness HV10 is generally not more than 2200) and surface finish are not ideal. In order to improve its hardness and surface finish, cemented carbide without a binder phase has been developed, which refers to a tungsten carbide-based cemented carbide with a binder phase metal (Co, Ni or Fe) content of less than 1wt.%. Since the binder phase content of this type of cemented carbide is extremely low, it is also called a tungsten carbide-based ceramic material. In this material system, the binder phase should be strictly referred to as a sintering aid. Compared with traditional cemented carbide, cemented carbide without a binder phase has excellent wear resistance, oxidation resistance, corrosion resistance and polishing properties. Therefore, it is widely used in wear-resistant fields such as ultra-mirror glass lens molds, water jet sand tubes, ultra-high pressure nozzles and high wear-resistant sealing rings. However, also due to the extremely low binder phase content, although its hardness HV10 is generally higher than 2300, its toughness is poor (usually K 1C =2MPa·m 1 / 2 ~3MPa·m 1 / 2 Even if the WC grain size is reduced to a submicron level of 0.4 μm, its toughness can only reach K 1C =4MPa·m 1 / 2 ~5MPa·m 1 / 2 This performance defect seriously limits its application range and makes it difficult to meet the application environment with impact loads, especially in the application of cutting tools.

[0005] Therefore, how to improve the toughness of binderless cemented carbide is an urgent problem to be solved in this field.

[0006] In view of this, the present invention is proposed. Summary of the Invention

[0007] The object of the present invention is to provide a binder-free cemented carbide and a preparation method thereof, aiming to significantly improve the toughness of the material without significantly reducing the hardness and causing a decrease in wear resistance.

[0008] Another object of the present invention is to provide the use of the above-mentioned binder-free cemented carbide in the preparation of cutting tools, molds or mining tools.

[0009] The present invention is achieved in that:

[0010] In a first aspect, the present invention provides a binderless cemented carbide, which is mainly composed of a hard phase; the grain morphology of the hard phase includes ultrafine crystals and lamellar crystals, the average particle size of the ultrafine crystals is 0.2μm to 0.4μm; the lamellar crystals are polygonal flakes, and the ratio of the polygonal equivalent circle diameter to the thickness of the lamellar crystals is 10 to 60:1.

[0011] In an optional embodiment, in the binderless cemented carbide, the volume content of the platelet crystals is 10% to 70%;

[0012] Preferably, the polygonal equivalent circle diameter of the plate-like crystal is 2 μm to 6 μm, and the thickness is 0.1 μm to 0.3 μm;

[0013] Preferably, the plate-like crystals are uniformly distributed in the polycrystalline matrix formed by ultrafine crystals.

[0014] In an optional embodiment, the chemical composition includes, by mass percentage, 0% to 1% binder phase metal, 0.01% to 0.03% boron, 0.5% to 3% titanium, 1% to 3% tantalum, 0% to 2% niobium, 0.8% to 1.2% chromium, 0.2% to 0.4% vanadium, 6.303% to 6.841% carbon, and the balance being tungsten and unavoidable impurities.

[0015] In an optional embodiment, the chemical composition includes, by mass percentage: 0.1% to 0.6% binder phase metal, 0.01% to 0.03% boron, 1% to 2% titanium, 1.5% to 2.5% tantalum, 0.5% to 1.5% niobium, 0.9% to 1.1% chromium, 0.2% to 0.4% vanadium, 6.423% to 6.653% carbon, and the balance is tungsten and unavoidable impurities;

[0016] Preferably, the binder phase metal is selected from at least one of cobalt, nickel and iron.

[0017] In a second aspect, the present invention provides a method for preparing a binderless cemented carbide according to any one of the aforementioned embodiments, comprising: preparing ingredients according to the chemical composition of the binderless cemented carbide, and mixing the raw materials to obtain a powdered mixture;

[0018] The powdered mixture is subjected to a sintering pre-treatment to obtain a green body to be sintered, and the green body to be sintered is sintered.

[0019] In an optional embodiment, the preparation process of the powdered mixture includes: mixing raw material powders weighed according to chemical composition by wet ball milling, and then drying;

[0020] Preferably, the raw material powder includes: binder phase metal powder, titanium tungsten carbide powder, tantalum carbide powder, niobium carbide powder, tungsten boride powder, chromium carbide powder, vanadium carbide powder, tungsten carbide powder and inevitable impurities;

[0021] Preferably, the binder phase metal powder is a single metal powder or alloy powder. When the binder phase metal powder is cobalt powder, the content of the close-packed hexagonal crystal structure in the powder is greater than 90 wt.%;

[0022] Preferably, the titanium tungsten carbide powder is a solid solution powder, and its chemical formula is (Ti x W 1-x )C, x=0.6~0.8;

[0023] Preferably, the tungsten boride powder is WB powder;

[0024] Preferably, the average particle size of the raw material powder as a whole is 0.1 μm to 1.5 μm; more preferably, the average particle size of the tungsten carbide powder is 0.1 μm to 0.4 μm, the average particle size of the binder metal powder is 0.2 μm to 0.6 μm, and the average particle size of the other raw material powders is 0.1 μm to 1.5 μm. The morphology of the raw material powder can be spherical, nearly spherical, or equiaxed.

[0025] In an optional embodiment, the mixing process using wet ball milling includes: first ball milling the other powders except the binder metal powder and the forming agent powder for 24 hours to 36 hours, then adding the binder metal powder and ball milling for 24 hours to 36 hours, and then adding the forming agent powder and ball milling for 15 hours to 20 hours;

[0026] Preferably, the forming agent powder is at least one of paraffin wax and polyethylene glycol;

[0027] Preferably, the ball-to-material mass ratio used in wet ball milling is (4-6):1;

[0028] Preferably, the medium used in wet ball milling is alcohol.

[0029] In an optional embodiment, the pre-sintering treatment process includes: shaping the powdered mixture into a primary body, and degreasing the primary body.

[0030] In an optional embodiment, the process of sintering the green body to be sintered includes: first vacuum sintering the green body to be sintered to obtain a pre-sintered body, and then hot isostatic pressing the pre-sintered body;

[0031] Preferably, the vacuum sintering temperature is 1540° C. to 1580° C.; the holding time is 1 h to 3 h, and the pressure is less than or equal to 30 Pa;

[0032] Preferably, the hot isostatic pressing sintering temperature is 1400° C. to 1450° C., the holding time is 1 h to 3 h, and the pressure is 150 MPa to 300 MPa;

[0033] More preferably, the hot isostatic pressing sintering is performed under an inert atmosphere.

[0034] In a third aspect, the present invention provides use of any of the binderless cemented carbides described in the aforementioned embodiments or any of the binderless cemented carbides prepared by the aforementioned preparation methods in the preparation of cutting tools, molds, or mining tools.

[0035] The present invention has the following beneficial effects: since the content of the binder phase metal in the binder-free cemented carbide provided by the present invention is less than or equal to 1 wt.%, it is difficult for WC grains to grow into a plate-like morphology by dissolving and precipitating in the binder phase metal. The inventors add elements such as boron, titanium, chromium, and vanadium and control their amounts to achieve directional control of the diffusion growth rate of each crystal plane and grain boundary of the WC grains, thereby controlling the growth of some grains to form an ultra-thin lamellae morphology. This achieves the preparation of polygonal lamellae-like morphology grains in the absence of a substantial amount of binder phase metal and without the need for flattening the raw material powder, effectively improving the toughness of the binder-free cemented carbide without significantly reducing its hardness and thereby degrading its wear resistance. BRIEF DESCRIPTION OF THE DRAWINGS

[0036] In order to more clearly illustrate the technical solutions of the embodiments of the present invention, the following briefly introduces the drawings required for use in the embodiments. It should be understood that the following drawings only illustrate certain embodiments of the present invention and therefore should not be regarded as limiting the scope. For ordinary technicians in this field, other relevant drawings can be obtained based on these drawings without paying any creative work.

[0037] Figure 1 This is a grain morphology image (low magnification image) of the fracture of the binderless cemented carbide prepared in Example 1.2 of the present invention;

[0038] Figure 2 This is a grain morphology image (high magnification image) of the fracture of the binderless cemented carbide prepared in Example 1.2 of the present invention;

[0039] Figure 3 This is a grain morphology image of the fracture of a conventional cemented carbide prepared in Comparative Example 1.1 provided by the present invention;

[0040] Figure 4 This is a grain morphology diagram of the fracture of the binderless cemented carbide prepared in Comparative Example 3.2 provided by the present invention. DETAILED DESCRIPTION

[0041] To make the purpose, technical solutions and advantages of the embodiments of the present invention clearer, the technical solutions in the embodiments of the present invention are described clearly and completely below. Where specific conditions are not specified in the embodiments, conventional conditions or conditions recommended by the manufacturer are used. Where the manufacturer of the reagents or instruments is not specified, all are conventional products that can be purchased commercially.

[0042] In an embodiment of the present invention, a binder-free cemented carbide is provided, whose chemical composition, calculated by mass percentage, includes: 0% to 1% binder phase metal, 0.01% to 0.03% boron, 0.5% to 3% titanium, 1% to 3% tantalum, 0% to 2% niobium, 0.8% to 1.2% chromium, 0.2% to 0.4% vanadium, 6.303% to 6.841% carbon, and the balance is tungsten and unavoidable impurities.

[0043] It should be noted that the content of the binder phase metal added to the raw materials of the binder-free cemented carbide provided in the embodiments of the present invention is less than or equal to 1 wt.%, and it is difficult for WC grains to grow into a plate-like morphology by dissolving and precipitating in the binder phase metal. The inventors add elements such as boron, titanium, chromium and vanadium (added in the form of carbide powder during the preparation process) and control their dosage to achieve directional control of the diffusion growth rate of each crystal plane and grain boundary of the WC grains, thereby controlling the growth of some grains to form a polygonal flake morphology. Polygonal flake morphology grains can be prepared under conditions of basically no binder phase metal and without the need for flattening the raw material powder, effectively improving the toughness of the binder-free cemented carbide without significantly reducing its hardness and thus reducing its wear resistance.

[0044] Specifically, the content of the binder phase metal can be 0.0%, 0.1%, 0.2%, 0.3%, 0.4%, 0.5%, 0.6%, 0.7%, 0.8%, 0.9%, 1.0%, etc.; the content of boron can be 0.01%, 0.02%, 0.03%, etc.; the content of titanium can be 0.5%, 0.8%, 1.0%, 1.2%, 1.5%, 1.7%, 2.0%, 2.2%, 2.5%, 2.7%, 3.0%, etc.; the content of tantalum can be 1.0%, 1.2%, 1.5%, 1.7%, 2.0%, 2.2%, 2.5%, 2.7%, 3.0%, etc. %, 2.2%, 2.5%, 2.7%, 3.0% and the like; the niobium content may be 0.0%, 0.2%, 0.5%, 0.8%, 1.0%, 1.2%, 1.5%, 1.7%, 2.0%; the chromium content may be 0.8%, 0.9%, 1.0%, 1.1%, 1.2% and the like; the vanadium content may be 0.2%, 0.3%, 0.4% and the like; the carbon content may be 6.303%, 6.4%, 6.5%, 6.6%, 6.7%, 6.8%, 6.841% and the like; the balance may be tungsten and unavoidable impurities.

[0045] In some embodiments, the chemical composition, measured by mass percentage, includes: binder metal 0.1%-0.6%, boron 0.01%-0.03%, titanium 1%-2%, tantalum 1.5%-2.5%, niobium 0.5%-1.5%, chromium 0.9%-1.1%, vanadium 0.2%-0.4%, carbon 6.423%-6.653%, and the balance being tungsten and unavoidable impurities. By optimizing the content of each element, the toughness and hardness of the alloy can be further improved.

[0046] In some embodiments, the binder phase metal is selected from at least one of cobalt, nickel, and iron, and may be any one or more of these. Binder-free cemented carbide is primarily composed of a hard phase, formed by in-situ reaction of the initial binder phase metal or alloy components with other components. "Mainly" here means essentially all hard phase, specifically, the hard phase accounts for at least 99% by mass.

[0047] In some embodiments, the grain morphology of the hard phase in the binderless cemented carbide includes ultrafine crystals and lamellar crystals, and may also include other grain morphologies. Alternatively, the grain morphology of all hard phases may be composed of both ultrafine crystals and lamellar crystals, with the lamellar crystals evenly distributed within a polycrystalline matrix formed by the ultrafine crystals. The volume fraction of the lamellar crystals in the cemented carbide grain morphology may be 10% to 70%, such as 10%, 20%, 30%, 40%, 50%, 60%, 70%, etc. The presence of the lamellar crystals significantly increases the fracture toughness of the alloy.

[0048] In some embodiments, the average grain size of the ultrafine crystals in the polycrystalline matrix is ​​0.2 μm to 0.4 μm, such as 0.2 μm, 0.3 μm, 0.4 μm, etc.

[0049] In some embodiments, the flake crystals are in the form of polygonal thin flakes, the polygonal equivalent circle diameter of the flake crystals is 2 μm to 6 μm, the thickness is 0.1 μm to 0.3 μm, and the ratio of the polygonal equivalent circle diameter to the thickness of the flake crystals is 10 to 60:1. The polygonal equivalent circle diameter is more than 10 times the thickness, and the flake crystals are very thin. Specifically, the polygonal equivalent circle diameter of the flake crystal can be 2μm, 3μm, 4μm, 5μm, 6μm, etc., the thickness can be 0.1μm, 0.2μm, 0.3μm, etc., and the ratio of the polygonal equivalent circle diameter to the thickness can be 10:1, 12:1, 15:1, 17:1, 20:1, 22:1, 25:1, 27:1, 30:1, 32:1, 35:1, 37:1, 40:1, 42:1, 45:1, 47:1, 50:1, 52:1, 55:1, 57:1, 60:1, etc.

[0050] The present invention also provides a method for preparing a binderless cemented carbide, comprising: preparing the binderless cemented carbide according to its chemical composition, so that the chemical composition of the prepared binderless cemented carbide meets the requirements for the elemental composition of the binderless cemented carbide. Specifically, the method may include the following steps:

[0051] S1. Mixing

[0052] The ingredients are prepared according to the chemical composition of the cemented carbide without a binder phase, and the powders of the raw materials are mixed to obtain a powdery mixture for standby use.

[0053] In some embodiments, the preparation process of the powdered mixture includes: mixing raw material powders weighed according to chemical composition by wet ball milling, and then drying. The drying method is not limited, and spray drying can be used.

[0054] In some embodiments, the raw material powder includes: binder phase metal powder, titanium tungsten carbide powder, tantalum carbide powder, niobium carbide powder, tungsten boride powder, chromium carbide powder, vanadium carbide powder, tungsten carbide powder, and inevitable impurities. The amount of raw materials used is proportioned according to the chemical composition of the target product.

[0055] It should be noted that chromium carbide powder and vanadium carbide powder help inhibit the grain boundary diffusion growth of WC grains. Compared with traditional cemented carbide, due to the extremely low binder phase content, the ultra-hard boride formed further reduces the dissolution of WC in the binder phase metal. The probability of WC grains growing through the dissolution-precipitation mechanism in the binder phase metal is very low. Most of the time, the grain refinement is achieved by inhibiting the grain boundary diffusion of WC grains by chromium carbide or vanadium carbide particles, thus forming a polycrystalline matrix with ultrafine grains with an average grain size of 0.2μm to 0.4μm. At the same time, the interaction between chromium and vanadium elements and boron elements will also affect the morphological characteristics of the platelet crystals and the uniformity of element distribution.

[0056] Furthermore, the bonding phase metal powder can be a single metal powder or an alloy powder. The single metal powder can be a single cobalt powder, nickel powder, iron powder, etc. The alloy powder can be an alloy raw material formed by two or three elements of cobalt, nickel, and iron.

[0057] In some embodiments, when the binder phase metal powder is cobalt powder, the content of hexagonal close-packed crystal structure (HCP) in the powder is greater than 90 wt.%. The cobalt powder with HCP structure is more brittle, which helps to slow down the cold welding agglomeration between particles during ball milling, thereby improving the uniformity of ball milling mixing.

[0058] In some embodiments, the titanium tungsten carbide powder is a solid solution powder, and its chemical formula is (Ti x W 1-x)C, x = 0.6 to 0.8, such as 0.6, 0.7, 0.8, etc. The use of titanium carbide tungsten solid solution powder helps to reduce the sintering temperature. At the same time, Ti atoms are easily deposited on the (0001) plane of the WC crystal, thereby hindering the diffusion growth of the (0001) plane of the WC grain; that is, it can inhibit the growth of WC grains along the

[0001] direction, thereby promoting the formation of lamellar crystals and controlling the content of lamellar crystals. In addition, compared with the use of TiC powder, the use of titanium carbide tungsten solid solution powder helps to reduce the sintering temperature, which is not only more conducive to the formation of lamellar crystals, but also conducive to sintering densification.

[0059] In some embodiments, the tungsten boride powder can be WB powder, which is a commercially available raw material. The density of WB powder is similar to that of WC powder, which helps to mix evenly during ball milling. At the same time, the presence of boron makes WB easy to react with the binder metal Co, Ni or Fe in situ to form WC powder. a Co b B c or W a Ni b B c or W a Fe b B c The hard phase of superhard boride can not only reduce the sintering temperature, but also help to improve the problem of reduced wear resistance caused by the presence of soft bonding phase metal alloy, thereby improving the red hardness of the material as a cutting tool; on the other hand, the formation of superhard boride phase can also improve the high-temperature oxidation resistance of the material; thirdly, and more importantly, the presence of boron element contributes to the formation of lamellar crystals, which may be because the above-mentioned superhard boride is easy to deposit on the (0001) surface of WC grains. Since Ti or Cr elements are strong boride-forming elements, they can also easily attract Ti atoms and Cr atoms to deposit on the (0001) surface of WC grains, hindering the diffusion growth of WC grains along the (0001) crystal plane to form columnar crystals, and can also accelerate the lateral growth rate of each columnar surface of WC grains to form a lamellar crystal morphology.

[0060] In some embodiments, the average particle size of the raw material powder as a whole is 0.1 μm to 1.5 μm, and the average particle size of all the above raw materials as a whole can be 0.1 μm, 0.5 μm, 0.7 μm, 1.0 μm, 1.2 μm, 1.5 μm, etc. Among them, the average particle size of tungsten carbide powder is 0.1 μm to 0.4 μm, the average particle size of binder metal powder is 0.2 μm to 0.6 μm, and the average particle size of other raw material powders is 0.1 μm to 1.5 μm. Other raw material powders refer to other types of powders other than tungsten carbide powder and binder metal powder. Specifically, the average particle size of the tungsten carbide powder can be 0.1 μm, 0.2 μm, 0.3 μm, 0.4 μm, etc., the average particle size of the binder metal powder can be 0.2 μm, 0.3 μm, 0.4 μm, 0.5 μm, 0.6 μm, etc., and the average particle size of other raw material powders can be 0.1 μm, 0.5 μm, 0.8 μm, 1.0 μm, 1.2 μm, 1.5 μm, etc. The morphology of the raw material powder can be spherical, nearly spherical, or equiaxed.

[0061] In some embodiments, the wet ball milling process includes first ball milling the remaining powders (i.e., the hard phase powder) except the binder metal powder and the forming agent powder for 24 to 36 hours, then adding the binder metal powder and ball milling for 24 to 36 hours, and then adding the forming agent powder and ball milling for 15 to 20 hours. The stepwise addition of the raw material powders during the rolling wet ball milling process achieves an ultrafine particle size, thereby increasing the activity of the raw material powders and improving the mixing uniformity of the raw material powders.

[0062] It should be noted that first, the other powders except the binder metal powder and the forming agent powder are ball-milled and mixed, so as to avoid the influence of the binder metal powder with good plasticity on the ball-milling crushing of the brittle hard phase powder, thereby improving the crushing efficiency and mixing uniformity, so as to promote the formation of highly active ultrafine particles, help to control the growth of lamellar crystals during subsequent sintering, and also help to form an ultrafine grain structure. Finally, the forming agent powder is added and mixed to avoid the influence of the forming agent powder on the ball milling process. Since the forming agent powder has good viscosity and plasticity, if it is added first, it will not only cause the raw material powder to agglomerate and be easily mixed unevenly, but also easily cause the forming agent itself to agglomerate and be mixed unevenly. Therefore, in the embodiment of the present invention, the hard phase powder is first crushed and mixed evenly, then mixed evenly with the binder metal powder, and finally mixed with the forming agent powder.

[0063] Specifically, the ball milling time of the hard phase powder can be 24h, 26h, 28h, 30h, 32h, 34h, 36h, etc., the ball milling time after adding the binding phase metal powder can be 24h, 26h, 28h, 30h, 32h, 34h, 36h, etc., and the ball milling time after adding the forming agent powder can be 15h, 16h, 17h, 18h, 19h, 20h, etc.

[0064] In some embodiments, the molding agent powder is at least one of paraffin wax and polyethylene glycol, and can be any one or a mixture of two of the above.

[0065] In some embodiments, the ball-to-material mass ratio used in wet ball milling is (4-6):1, such as 4:1, 5:1, 6:1, etc.; the medium used in wet ball milling can be but is not limited to alcohol.

[0066] S2. Pre-sintering treatment

[0067] The powdered mixed material is subjected to a pre-sintering treatment to obtain a green body to be sintered. The specific process of the pre-sintering treatment is not limited.

[0068] In some embodiments, the pre-sintering treatment process includes: shaping the powdered mixture into a green body, and degreasing the green body to remove the forming agent in the green body.

[0069] S3, sintering

[0070] The green body to be sintered can be sintered by combining vacuum sintering and hot isostatic pressing, which is beneficial for improving the toughness and hardness of the prepared binderless cemented carbide. In actual operation, the process of sintering the green body to be sintered includes: first vacuum sintering the green body to be sintered to obtain a pre-sintered body, and then hot isostatic pressing the pre-sintered body to obtain the final dense binderless cemented carbide.

[0071] In some embodiments, the vacuum sintering temperature is 1540°C to 1580°C; the holding time is 1 hour to 3 hours, and the pressure is less than or equal to 30 Pa. By selecting an appropriate vacuum sintering temperature, it helps promote sintering densification while preventing abnormal grain growth. If the vacuum sintering temperature is lower than 1540°C, the pre-sintered body contains a large number of open pores and cannot meet the density requirements of the pre-sintered body for the next step of hot isostatic pressing. If the vacuum sintering temperature is higher than 1580°C, the side length and thickness of the lamellar crystals will increase, and the ultrafine grains will also grow, resulting in a decrease in the toughness and hardness of the binderless cemented carbide.

[0072] Specifically, the vacuum sintering temperature can be 1540°C, 1550°C, 1560°C, 1570°C, 1580°C, etc., and the holding time can be 1h, 2h, 3h, etc.

[0073] In some embodiments, the hot isostatic pressing (HIP) sintering temperature is 1400°C to 1450°C, the holding time is 1 hour to 3 hours, and the pressure is 150 MPa to 300 MPa. The HIP sintering is performed in an inert atmosphere (e.g., argon). By optimizing the HIP sintering parameters to fully ensure that the grains in the pre-sintered body no longer grow significantly, the residual porosity in the pre-sintered body is eliminated, thereby obtaining a dense, binder-free cemented carbide.

[0074] Specifically, the hot isostatic pressing sintering temperature can be 1400°C, 1410°C, 1420°C, 1430°C, 1440°C, 1450°C, etc., the holding time can be 1h, 2h, 3h, etc., and the pressure can be 150MPa, 180MPa, 200MPa, 220MPa, 250MPa, 270MPa, 300MPa, etc. Generally, when the hot isostatic pressing temperature is slightly higher, the sintering pressure can be appropriately reduced; when the hot isostatic pressing temperature is slightly lower, the sintering pressure can be appropriately increased. Usually, in order to avoid grain growth, a lower hot isostatic pressing temperature and a higher sintering pressure are preferred.

[0075] It should be noted that the binderless cemented carbide and its preparation method provided in the embodiment of the present invention can prepare a binderless cemented carbide composed of ultrafine crystals and lamellar crystals, and the lamellar crystals are evenly distributed in the polycrystalline matrix composed of ultrafine crystals. Due to the presence of lamellar crystals, its fracture toughness is greatly increased. The preparation method provided in the embodiment of the present invention does not require flattening of the raw material powder during the preparation process, nor does it require growth through the dissolution-precipitation mechanism of WC in the binder phase metal to form lamellar crystals. The preparation method is simple and effective. In addition, the binderless cemented carbide prepared in the embodiment of the present invention retains the polycrystalline matrix of ultrafine crystals, and weakens the influence of the softer bonding pure metal phase or bonding metal alloy phase, and is mainly composed of a hard phase. Therefore, a binderless cemented carbide with high toughness and hardness is obtained.

[0076] The embodiments of the present invention also provide the use of the above-mentioned binderless cemented carbide in the preparation of cutting tools, molds, or mining tools, utilizing its high toughness and hardness to make the prepared cutting tools, molds, or mining tools and other products also have excellent performance.

[0077] The following is a further detailed description of the structural characteristics and properties of the binderless cemented carbide prepared by the present invention in conjunction with the examples.

[0078] Experimental group 1

[0079] Experimental Group 1 includes Examples 1.1 to 1.5 and Comparative Examples 1.1 to 1.2, and mainly adjusts the Co content, as follows:

[0080] Example 1.1

[0081] This embodiment provides a method for preparing a binderless cemented carbide, which is prepared by the following steps:

[0082] (1) The raw material powders were accurately weighed according to the proportion, and mixed by rolling wet ball milling with stepwise addition of raw material powders, with a ball-to-material ratio of 5:1, and the medium used in the wet ball milling was alcohol; and the mixed material powder was prepared by spray drying. Wherein, the raw material powders were composed of: (Ti 0.75 W 0.25 )C 3.922%, TaC 2.133%, NbC 1.129%, WB 0.360%, Cr3C2 1.154%, VC 0.371%, the balance is WC. The WC, (Ti 0.75 W 0.25 The morphologies of the raw material powders of )C, TaC, NbC, WB, Cr3C2, and VC are all equiaxed. The step-by-step process of adding the raw material powders includes: first, ball milling all brittle powders except the binder metal cobalt powder and the molding agent for 30 hours; then, adding the binder metal cobalt powder and ball milling for 30 hours (this step is omitted if the initial binder metal powder content is zero); finally, adding the paraffin molding agent powder and ball milling for 18 hours.

[0083] (2) The blank is pressed by molding.

[0084] (3) Degreasing the obtained green body is performed using an argon negative pressure process.

[0085] (4) The degreased green body is vacuum sintered to obtain a pre-sintered body, the sintering temperature is 1560°C, the holding time is 2h, and the pressure is less than or equal to 30Pa.

[0086] (5) The pre-sintered body is subjected to hot isostatic pressing at a sintering temperature of 1430°C, a holding time of 2 h, and an argon pressure of 200 MPa.

[0087] Example 1.2-Example 1.5

[0088] The only difference between Examples 1.2 to 1.5 and Example 1.1 is that the composition of the raw material powder is adjusted for the cemented carbide without a binder phase. See Table 1 for details.

[0089] Comparative Example 1.1-Comparative Example 1.2

[0090] Comparative Examples 1.1 and 1.2 differ from Example 1 only in that the composition of the raw material powder and the sintering process are adjusted for the cemented carbide. The composition of the raw material powder in the cemented carbide is shown in Table 1. The sintering processes used in Comparative Examples 1.1 and 1.2 are as follows:

[0091] The vacuum sintering temperatures used in Comparative Examples 1.1 and 1.2 were 1420°C and 1400°C, respectively, the holding time was 2h, and the pressure was less than or equal to 30Pa; then a low-pressure sintering process was adopted, and the sintering temperatures of Comparative Examples 1.1 and 1.2 were 1400°C and 1380°C, respectively, the holding time was 2h, and the argon pressure was 5MPa.

[0092] Note: Comparative Examples 1.1 and 1.2 are based on the formula composition of traditional cemented carbide with a high cobalt content. If the same vacuum sintering temperature (e.g., 1560°C) as that of binderless cemented carbide is used, severe over-burning will result.

[0093] Table 1 Chemical formula and ratio of raw material powder used in test group 1 (wt.%)

[0094]

[0095]

[0096] The chemical composition of the prepared cemented carbide was tested using the following methods: Metal element content was determined according to the national standard GB / T26050-2010, "Determination of Metal Element Content in Cemented Carbide by X-ray Red Fluorescence Fusion Method"; total carbon content was determined according to the national standard GB / T42275-2022, "Determination of Total Carbon in Cemented Carbide by High-Frequency Combustion Infrared Absorption Method / Thermal Conductivity Method"; and boron content was determined according to the national standard GB / T39138.3-2020, "Chemical Analysis Methods for Gold-Nickel-Chromium-Fe-Silicon-Boron Alloys - Part 3: Determination of Chromium, Iron, Silicon, and Boron Contents - Inductively Coupled Plasma Optical Emission Spectrometry." The test results are shown in Table 2.

[0097] Table 2 Chemical composition ratio of cemented carbide in test group 1 (wt.%)

[0098] Case No. Co Ti Ta Nb B Cr V W C Example 1.1 0.0 1.50 2.00 1.00 0.02 1.00 0.30 87.620 6.560 Example 1.2 0.2 1.50 2.00 1.00 0.02 1.00 0.30 87.433 6.547 Example 1.3 0.5 1.50 2.00 1.00 0.02 1.00 0.30 87.151 6.529 Example 1.4 0.8 1.50 2.00 1.00 0.02 1.00 0.30 86.869 6.511 Example 1.5 1.0 1.50 2.00 1.00 0.02 1.00 0.30 86.682 6.498 Comparative Example 1.1 6.0 1.50 2.00 1.00 0.02 1.00 0.30 81.988 6.192 Comparative Example 1.2 9.0 1.50 2.00 1.00 0.02 1.00 0.30 79.172 6.008

[0099] The microstructure and properties of the cemented carbides obtained by testing Examples 1.1 to 1.5 and Comparative Examples 1.1 to 1.2 are shown in Table 3.

[0100] Test method:

[0101] (1) Test of the morphology and size of flake crystals: The test was conducted by SEM observation, specifically as follows: Observe the morphology characteristics by SEM; randomly test at least 500 flake crystals for each sample to obtain the thickness and polygonal area of ​​each flake crystal. The polygon was equivalent to a circle using the equal area method, and the equivalent circle diameter was obtained by the following formula:

[0102] D i =(4S i / π)1 / 2

[0103] Where S i is the measured value of the polygonal area of ​​the i-th platelet; D i is the polygonal equivalent circle diameter of the i-th platelet.

[0104] (2) Test of the flake crystal content: Based on the polygonal area and thickness of the flake crystals obtained in the test, the volume of a single flake crystal is the product of the polygonal area and the thickness; then, based on the number of flake crystals in the obtained test volume, the volume content of the flake crystals is calculated.

[0105] (3) Test of ultrafine grain size of matrix: The test was conducted by SEM observation. The test method was based on the national standard GB / T3488.2-2018 “Metallographic determination of cemented carbide microstructure Part 2: Measurement of WC grain size”.

[0106] (4) The fracture toughness and hardness tests were carried out in accordance with the industry standard JB / T12616-2016 “Test Method for Fracture Toughness of Cemented Carbide Tool Matrix Material” and the national standard GB / T7997-2014 “Test Method for Vickers Hardness of Cemented Carbide”.

[0107] Table 3 Comparison of the microstructure and performance parameters of cemented carbide obtained in test group 1

[0108]

[0109] As shown in Table 3, the fracture toughness of the binderless cemented carbides prepared in Examples 1.1 to 1.5 is higher than 7 MPa·m 1 / 2 , and the hardness is higher than 2400. In Examples 1.1 to 1.5, as the cobalt content in the binder phase increases, the atomic diffusion capacity is enhanced, the polygonal equivalent circle diameter of the formed plate-like grains increases, the thickness also increases, the ratio of the polygonal equivalent circle diameter to the thickness decreases, and the content of plate-like grains also increases slightly; accordingly, the size of the ultrafine grains in the matrix also increases slightly, and the typical organizational structure is shown in FIG. Figure 1 and Figure 2 shown.

[0110] The fracture toughness of the cemented carbides prepared in Comparative Examples 1.1 and 1.2 is also higher than 7 MPa·m 1 / 2 , but the hardness is only below 2108. Comparative Examples 1.1 and 1.2 have high Co contents. Due to the enhanced dissolution and precipitation of WC in the liquid phase cobalt during sintering, in the absence of flattened raw material powder particles as crystal nuclei, even in the presence of boron, conventional triangular prism-shaped grains are formed instead of lamellar crystals. Figure 3This is the grain morphology of the cemented carbide fracture surface obtained in Comparative Example 1.1. On the one hand, the increased cobalt content in the binder phase increases the fracture toughness of the cemented carbide while decreasing its hardness. On the other hand, as the cobalt content in the binder phase increases, the morphology of the lamellar crystals shifts from flaky to columnar, which in turn decreases the fracture toughness of the alloy. Therefore, the comprehensive results show that when the cobalt content in the binder phase reaches 6%, the toughness of Comparative Examples 1.1 and 1.2 does not increase significantly compared to Example 1.5, but the hardness decreases significantly, by 9% to 16%.

[0111] Experimental Group 2

[0112] Experimental Group 2 includes Example 1.2, Example 2.1-Example 2.2, and Comparative Example 2.1-Comparative Example 2.2. The main adjustment is the content of the boron element, and the type and ratio of the raw material powder are adjusted accordingly, as shown in Table 4:

[0113] Table 4 Chemical formula and ratio of raw material powder used in test group 2 (wt.%)

[0114] Case No. Co <![CDATA[(Ti 0.75 W 0.25 )C]]> TaC NbC WB <![CDATA[Cr3C2]]> VC WC Comparative Example 2.1 0.2 3.922 2.133 1.129 0.000 1.154 0.371 margin Example 2.1 0.2 3.922 2.133 1.129 0.180 1.154 0.371 margin Example 1.2 0.2 3.922 2.133 1.129 0.360 1.154 0.371 margin Example 2.2 0.2 3.922 2.133 1.129 0.540 1.154 0.371 margin Comparative Example 2.2 0.2 3.922 2.133 1.129 0.720 1.154 0.371 margin

[0115] The preparation methods of Examples 2.1-2.2 and Comparative Examples 2.1-2.2 refer to Example 1.1.

[0116] The chemical composition test results of the binder-free cemented carbides obtained from Examples 2.1, 1.2, 2.2 and Comparative Examples 2.1-2.2 are shown in Table 5.

[0117] Table 5 Chemical composition ratio of cemented carbide without binder phase in test group 2 (wt.%)

[0118] Case No. Co Ti Ta Nb B Cr V W C Comparative Example 2.1 0.2 1.50 2.00 1.00 0.00 1.00 0.30 87.431 6.569 Example 2.1 0.2 1.50 2.00 1.00 0.01 1.00 0.30 87.432 6.558 Example 1.2 0.2 1.50 2.00 1.00 0.02 1.00 0.30 87.433 6.547 Example 2.2 0.2 1.50 2.00 1.00 0.03 1.00 0.30 87.434 6.536 Comparative Example 2.2 0.2 1.50 2.00 1.00 0.04 1.00 0.30 87.435 6.525

[0119] The microstructure and properties of the binder-free cemented carbide obtained by testing Example 1.2, Example 2.1-Example 2.2, and Comparative Example 2.1-Comparative Example 2.2 are shown in Table 6.

[0120] Table 6 Comparison of the microstructure and performance parameters of the binderless cemented carbide obtained in test group 2

[0121]

[0122] As shown in Table 6, when no boron is added, no flake grains are formed, and an ultrafine-grained cemented carbide structure without a binder phase is obtained, which has a high hardness but low toughness. When boron is added, the hardness decreases due to the formation of flake crystals, but the toughness increases. As the boron content increases, on the one hand, the ratio of the polygonal equivalent circle diameter to the thickness of the flake crystals increases significantly, the content of flake crystals also increases, and its toughness increases; on the other hand, due to the formation of W aCo b B c The superhard compound phase improves the low hardness of the binder phase metal cobalt, and its hardness is also increased. However, if the boron addition is too high, the grains tend to grow, that is, the size of the platelets and the size of the matrix grains increase, which will lead to a decrease in the fracture toughness and hardness of the alloy.

[0123] Experimental Group 3

[0124] Experimental Group 3 includes Example 1.2, Example 3.1-Example 3.4, and Comparative Examples 3.1-Comparative Examples 3.3. The main adjustment was the content of titanium element, and the type and ratio of the raw material powders were adjusted accordingly, as shown in Table 7:

[0125] Table 7 Chemical formula and ratio of raw material powder used in test group 3 (wt.%)

[0126] Case No. Co <![CDATA[(Ti 0.75 W 0.25 )C]]> TaC NbC WB <![CDATA[Cr3C2]]> VC WC Comparative Example 3.1 0.2 0.000 2.133 1.129 0.360 1.154 0.371 margin Comparative Example 3.2 0.2 0.523 2.133 1.129 0.360 1.154 0.371 margin Example 3.1 0.2 1.307 2.133 1.129 0.360 1.154 0.371 margin Example 3.2 0.2 2.615 2.133 1.129 0.360 1.154 0.371 margin Example 1.2 0.2 3.922 2.133 1.129 0.360 1.154 0.371 margin Example 3.3 0.2 5.229 2.133 1.129 0.360 1.154 0.371 margin Example 3.4 0.2 7.843 2.133 1.129 0.360 1.154 0.371 margin Comparative Example 3.3 0.2 10.458 2.133 1.129 0.360 1.154 0.371 margin

[0127] The preparation methods of Examples 3.1 to 3.4 and Comparative Examples 3.1 to 3.3 refer to Example 1.

[0128] Table 8 shows the chemical composition test results of the binderless cemented carbides of Example 1.2, Example 3.1 to Example 3.4, and Comparative Example 3.1 to Comparative Example 3.3.

[0129] Table 8 Chemical composition ratio (wt.%) of cemented carbide without binder phase in test group 3

[0130] Case No. Co Ti Ta Nb B Cr V W C Comparative Example 3.1 0.2 0.00 2.00 1.00 0.02 1.00 0.30 89.194 6.286 Comparative Example 3.2 0.2 0.20 2.00 1.00 0.02 1.00 0.30 88.959 6.321 Example 3.1 0.2 0.50 2.00 1.00 0.02 1.00 0.30 88.607 6.373 Example 3.2 0.2 1.00 2.00 1.00 0.02 1.00 0.30 88.020 6.460 Example 1.2 0.2 1.50 2.00 1.00 0.02 1.00 0.30 87.433 6.547 Example 3.3 0.2 2.00 2.00 1.00 0.02 1.00 0.30 86.846 6.634 Example 3.4 0.2 3.00 2.00 1.00 0.02 1.00 0.30 85.672 6.808 Comparative Example 3.3 0.2 4.00 2.00 1.00 0.02 1.00 0.30 84.498 6.982

[0131] The microstructure and properties of the binder-free cemented carbide were obtained by testing Example 1.2, Example 3.1 to Example 3.4, and Comparative Example 3.1 to Comparative Example 3.3, as shown in Table 9.

[0132] Table 9 Comparison of the microstructure and performance parameters of the binderless cemented carbide obtained in test group 3

[0133]

[0134] From Table 9, we can see that (Ti 0.75 W 0.25 ) As the content of C increases, the interaction between boron and titanium and tungsten elements is enhanced, the content of plate crystals increases, the toughness of the alloy increases, and the hardness of the alloy decreases. Figure 4 The grain morphology of the fracture surface of comparative example 3.2 is shown in FIG. When the content of the plate crystal is about 40 vol.% to 60 vol.%, and the average particle size of the matrix ultrafine crystal is 0.2 μm to 0.4 μm, the comprehensive performance of the alloy reaches the best. 0.75 W 0.25) C content continues to increase, on the one hand, the content of platelets will exceed 60vol.%, due to the bridge phenomenon between the platelets, resulting in a decrease in the fracture toughness of the alloy. 0.75 W 0.25 ) When the addition amount of C is too high, the grains will grow easily, the thickness of the lamellar crystals will increase, the size of the ultrafine crystals will increase and the volume content will decrease, resulting in a decrease in the hardness and toughness of the alloy.

[0135] Experimental Group 4

[0136] Experimental Group 4 includes Example 1.2, Example 4.1-Example 4.4, and Comparative Examples 4.1-Comparative Examples 4.3. The main adjustments were made to the content of tantalum or niobium elements, and the types and proportions of the raw material powders were adjusted accordingly. The details are shown in Table 10:

[0137] Table 10 Chemical formula and ratio of raw material powder used in test group 4 (wt.%)

[0138]

[0139]

[0140] The preparation methods of Examples 4.1 to 4.4 and Comparative Examples 4.1 to 4.3 refer to Example 1.

[0141] Table 11 shows the chemical composition test results of the binderless cemented carbides obtained from Example 1.2, Example 4.1 to Example 4.4, and Comparative Example 4.1 to Comparative Example 4.3.

[0142] Table 11 Chemical composition ratio (wt.%) of cemented carbide without binder phase in test group 4

[0143] Case No. Co Ti Ta Nb B Cr V W C Comparative Example 4.1 0.2 1.50 0.00 1.00 0.02 1.00 0.30 89.435 6.545 Example 4.1 0.2 1.50 1.00 1.00 0.02 1.00 0.30 88.434 6.546 Example 1.2 0.2 1.50 2.00 1.00 0.02 1.00 0.30 87.433 6.547 Example 4.2 0.2 1.50 3.00 1.00 0.02 1.00 0.30 86.432 6.548 Comparative Example 4.2 0.2 1.50 4.00 1.00 0.02 1.00 0.30 85.431 6.549 Example 4.3 0.2 1.50 2.00 0.00 0.02 1.00 0.30 88.493 6.487 Example 1.2 0.2 1.50 2.00 1.00 0.02 1.00 0.30 87.433 6.547 Example 4.4 0.2 1.50 2.00 2.00 0.02 1.00 0.30 86.373 6.607 Comparative Example 4.3 0.2 1.50 2.00 3.00 0.02 1.00 0.30 85.313 6.667

[0144] The microstructure and properties of the cemented carbides obtained by testing Example 1.2, Example 4.1 to Example 4.4, and Comparative Example 4.1 to Comparative Example 4.3 are shown in Table 12.

[0145] Table 12 Comparison of the microstructure and performance parameters of the binderless cemented carbide obtained in test group 4

[0146]

[0147]

[0148] Table 12 shows that changes in the TaC and NbC contents have little effect on the formation of platelets. However, as the TaC and NbC contents decrease, the WC content must be increased accordingly, resulting in an increase in the alloy's hardness and a slight decrease in fracture toughness. The primary purpose of adding TaC and NbC is to improve the alloy's oxidation resistance.

[0149] Experimental Group 5

[0150] Experimental Group 5 includes Example 1.2, Example 5.1 to Example 5.4, and Comparative Examples 5.1 to Comparative Examples 5.4. The main adjustments were made to the content of chromium or vanadium, and the types and ratios of the raw material powders were adjusted accordingly. The details are shown in Table 13:

[0151] Table 13 Chemical formula and ratio of raw material powder used in test group 5 (wt.%)

[0152] Case No. Co <![CDATA[(Ti 0.75 W 0.25 )C]]> TaC NbC WB <![CDATA[Cr3C2]]> VC WC Comparative Example 5.1 0.2 3.922 2.133 1.129 0.360 0.577 0.371 margin Example 5.1 0.2 3.922 2.133 1.129 0.360 0.923 0.371 margin Example 1.2 0.2 3.922 2.133 1.129 0.360 1.154 0.371 margin Example 5.2 0.2 3.922 2.133 1.129 0.360 1.385 0.371 margin Comparative Example 5.2 0.2 3.922 2.133 1.129 0.360 1.731 0.371 margin Comparative Example 5.3 0.2 3.922 2.133 1.129 0.360 1.154 0.124 margin Example 5.3 0.2 3.922 2.133 1.129 0.360 1.154 0.248 margin Example 1.2 0.2 3.922 2.133 1.129 0.360 1.154 0.371 margin Example 5.4 0.2 3.922 2.133 1.129 0.360 1.154 0.494 margin Comparative Example 5.4 0.2 3.922 2.133 1.129 0.360 1.154 0.618 margin

[0153] The preparation methods of Examples 5.1 to 5.4 and Comparative Examples 5.1 to 5.4 refer to Example 1.

[0154] Table 14 shows the chemical composition test results of the binderless cemented carbides obtained from Example 1.2, Examples 5.1 to 5.4, and Comparative Examples 5.1 to 5.4.

[0155] Table 14 Chemical composition ratio (wt.%) of cemented carbide without binder phase in test group 5

[0156] Case No. Co Ti Ta Nb B Cr V W C Comparative Example 5.1 0.2 1.50 2.00 1.00 0.02 0.50 0.30 87.974 6.506 Example 5.1 0.2 1.50 2.00 1.00 0.02 0.80 0.30 87.649 6.531 Example 1.2 0.2 1.50 2.00 1.00 0.02 1.00 0.30 87.433 6.547 Example 5.2 0.2 1.50 2.00 1.00 0.02 1.20 0.30 87.216 6.564 Comparative Example 5.2 0.2 1.50 2.00 1.00 0.02 1.50 0.30 86.891 6.589 Comparative Example 5.3 0.2 1.50 2.00 1.00 0.02 1.00 0.10 87.665 6.515 Example 5.3 0.2 1.50 2.00 1.00 0.02 1.00 0.20 87.548 6.532 Example 1.2 0.2 1.50 2.00 1.00 0.02 1.00 0.30 87.433 6.547 Example 5.4 0.2 1.50 2.00 1.00 0.02 1.00 0.40 87.317 6.563 Comparative Example 5.4 0.2 1.50 2.00 1.00 0.02 1.00 0.50 87.201 6.579

[0157] The microstructure and properties of the binder-free cemented carbide were obtained by testing Example 1.2, Example 5.1 to Example 5.4, and Comparative Example 5.1 to Comparative Example 5.4, as shown in Table 15.

[0158] Table 15 Comparison of the microstructure and performance parameters of the binderless cemented carbide obtained in test group 5

[0159]

[0160] Table 15 shows that when the Cr3C2 and VC contents are low, the growth inhibition of the matrix grains is weak, the matrix grains grow significantly, and the alloy's hardness is low. As the Cr3C2 and VC contents increase, the size of the lamellar crystals decreases, the ratio of the polygonal equivalent circle diameter to the thickness of the lamellar crystals decreases, and the matrix grain size also decreases, while the lamellar crystal content has little effect. Therefore, the alloy's hardness increases and its toughness decreases. However, when the Cr3C2 and VC contents are too high, agglomeration tends to occur during sintering. This is likely due to the presence of boron, which easily attracts Cr to agglomerate, thereby increasing the alloy's brittleness. The fracture toughness then shows a trend of increasing and then decreasing.

[0161] Experimental Group 6

[0162] This comparative example 6.1 provides a binderless cemented carbide and a preparation method thereof, which differs from Example 1.2 only in that: the (Ti 0.75 W 0.25 ) The C raw material powder is replaced by TiC powder, and the rest is the same as Example 1.2.

[0163] The microstructure and properties of the binder-free cemented carbides obtained in Comparative Example 6.1 and Example 1.2 are shown in Table 16.

[0164] Table 16 Comparison of properties and structural parameters of cemented carbide without binder phase

[0165]

[0166] As can be seen from Table 16, since the direct use of TiC powder will cause the temperature required for sintering to be dense to increase, the vacuum sintering temperature of Comparative Example 6.1 is relatively low, resulting in the residual pores in the pre-sintered body being mostly open pores, and the density requirement of the pre-sintered body for the next step of hot isostatic pressing is not met. Therefore, the material is not sintered densely and no further testing is carried out.

[0167] Experimental Group 7

[0168] This comparative example provides a cemented carbide and a preparation method thereof, which differs from Example 1.2 only in that the vacuum sintering temperature in Example 1.2 is lowered to 1520°C (Comparative Example 7.1) and increased to 1600°C (Comparative Example 7.2), respectively. The rest is the same as Example 1.2.

[0169] The microstructure and properties of the binder-free cemented carbides obtained in Comparative Examples 7.1 and 7.2 and Example 1.2 are shown in Table 17.

[0170] Table 17 Comparison of properties and microstructure parameters of the obtained binder-free cemented carbide

[0171]

[0172] As shown in Table 17, lowering the vacuum sintering temperature significantly increases the residual porosity in the pre-sintered body of Comparative Example 7.1, due to the relatively low vacuum sintering temperature. This pre-sintered body density requirement for hot isostatic pressing (HIP) was not met, resulting in a poorly sintered material and further testing was not performed. Increasing the vacuum sintering temperature significantly increases grain size, with the polygonal equivalent circle diameter / thickness ratio of the platelets in Comparative Example 7.2 decreasing dramatically. The matrix grains have grown to the micron level, ultimately leading to a sharp decrease in both fracture toughness and hardness.

[0173] The above are merely preferred embodiments of the present invention and are not intended to limit the present invention. Those skilled in the art will readily appreciate that various modifications and variations of the present invention are possible. Any modifications, equivalent substitutions, or improvements made within the spirit and principles of the present invention shall be included within the scope of protection of the present invention.

Claims

1. A binderless cemented carbide, characterized in that: The cemented carbide is mainly composed of a hard phase; the grain morphology of the hard phase includes ultrafine crystals and lamellar crystals, the average particle size of the ultrafine crystals is 0.2 μm to 0.4 μm; the lamellar crystals are polygonal flakes, and the ratio of the polygonal equivalent circle diameter to the thickness of the lamellar crystals is 10 to 60:1; Calculated by mass percentage, its chemical composition includes: binder phase metal 0% to 1%, boron 0.01% to 0.03%, titanium 0.5% to 3%, tantalum 1% to 3%, niobium 0% to 2%, chromium 0.8% to 1.2%, vanadium 0.2% to 0.4%, carbon 6.303% to 6.841%, and the balance is tungsten and unavoidable impurities.

2. The binder-free cemented carbide according to claim 1, characterized in that: In the binder-free cemented carbide, the volume content of the platelet crystals is 10% to 70%.

3. The binderless cemented carbide according to claim 1, characterized in that: The polygonal equivalent circle diameter of the plate-like crystal is 2 μm to 6 μm, and the thickness is 0.1 μm to 0.3 μm.

4. The binderless cemented carbide according to claim 1, characterized in that: The plate-like crystals are uniformly distributed in the polycrystalline matrix formed by the ultrafine crystals.

5. The binder-free cemented carbide according to claim 1, characterized in that: Calculated by mass percentage, its chemical composition includes: binder phase metal 0.1% to 0.6%, boron 0.01% to 0.03%, titanium 1% to 2%, tantalum 1.5% to 2.5%, niobium 0.5% to 1.5%, chromium 0.9% to 1.1%, vanadium 0.2% to 0.4%, carbon 6.423% to 6.653%, and the balance is tungsten and unavoidable impurities.

6. The binder-free cemented carbide according to claim 1, characterized in that: The binder phase metal is selected from at least one of cobalt, nickel and iron.

7. A method for preparing a binder-free cemented carbide according to any one of claims 1 to 6, characterized in that: include: The raw materials are mixed according to the chemical composition of the binderless cemented carbide to obtain a powdered mixture; The powdered mixed material is subjected to a pre-sintering treatment to obtain a green body to be sintered, and the green body to be sintered is sintered.

8. The preparation method according to claim 7, characterized in that The preparation process of the powdered mixture includes: mixing raw material powders weighed according to chemical composition by wet ball milling, and then drying.

9. The preparation method according to claim 7, characterized in that The raw material powder includes: binder phase metal powder, titanium tungsten carbide powder, tantalum carbide powder, niobium carbide powder, tungsten boride powder, chromium carbide powder, vanadium carbide powder, tungsten carbide powder and inevitable impurities.

10. The preparation method according to claim 9, characterized in that The binder phase metal powder is a single metal powder or alloy powder. When the binder phase metal powder is cobalt powder, the content of close-packed hexagonal crystal structure in the powder is greater than 90 wt.%.

11. The preparation method according to claim 9, characterized in that The titanium tungsten carbide powder is a solid solution powder, and its chemical formula is (Ti x W 1-x )C, x=0.6~0.

8.

12. The preparation method according to claim 9, characterized in that The tungsten boride powder is WB powder.

13. The preparation method according to claim 9, characterized in that The average particle size of the raw material powder as a whole is 0.1 μm to 1.5 μm.

14. The preparation method according to claim 13, characterized in that The average particle size of the tungsten carbide powder is 0.1 μm to 0.4 μm, the average particle size of the binder metal powder is 0.2 μm to 0.6 μm, and the average particle size of other raw material powders is 0.1 μm to 1.5 μm.

15. The preparation method according to claim 8, characterized in that The mixing process using wet ball milling includes: first ball milling the other powders except the binder phase metal powder and the forming agent powder for 24h to 36h, then adding the binder phase metal powder and ball milling for 24h to 36h, and then adding the forming agent powder and ball milling for 15h to 20h.

16. The preparation method according to claim 15, characterized in that The molding agent powder is at least one of paraffin wax and polyethylene glycol.

17. The preparation method according to claim 15, characterized in that The ball-to-material mass ratio used in the wet ball milling is (4-6):

1.

18. The preparation method according to claim 15, characterized in that The medium used in the wet ball milling is alcohol.

19. The preparation method according to claim 7, characterized in that: The pre-sintering treatment process includes: shaping the powdered mixture into a primary body, and degreasing the primary body.

20. The preparation method according to claim 7, characterized in that The process of sintering the green body to be sintered comprises: firstly vacuum sintering the green body to be sintered to obtain a pre-sintered body, and then hot isostatic pressing the pre-sintered body.

21. The preparation method according to claim 20, characterized in that The vacuum sintering temperature is 1540° C. to 1580° C.; the holding time is 1 hour to 3 hours; and the pressure is less than or equal to 30 Pa.

22. The preparation method according to claim 20, characterized in that The hot isostatic pressing sintering temperature is 1400° C. to 1450° C., the holding time is 1 hour to 3 hours, and the pressure is 150 MPa to 300 MPa.

23. The preparation method according to claim 22, characterized in that The hot isostatic pressing sintering is performed under an inert atmosphere.

24. Use of the binderless cemented carbide according to any one of claims 1 to 6 or the binderless cemented carbide prepared by the preparation method according to any one of claims 7 to 23 in the preparation of cutting tools, molds, or mining tools.

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