Binder-phase-free hard alloy, and preparation method therefor and use thereof

Abstract

The present invention relates to the technical field of hard alloys. Disclosed are a binder-phase-free hard alloy, and a preparation method therefor and the use thereof. Since the content of a binder-phase metal in the binder-phase-free hard alloy as provided is less than or equal to 1 wt%, WC crystal grains are difficult to grow into a plate-shaped morphology by means of dissolution-precipitation in the binder-phase metal, and the diffusion growth rate of each crystal plane and crystal boundary of the WC crystal grains is directionally controlled by adding elements such as boron, titanium, chromium and vanadium and controlling the amounts thereof; therefore, a portion of the crystal grains is controlled to grow into a polygonal lamellar morphology, crystal grains with a polygonal lamellar morphology are prepared under the conditions of substantially no binder-phase metal being present and flattening treatment not being required for raw material powder, the toughness of the binder-phase-free hard alloy is effectively improved, and the hardness is prevented from being significantly reduced to decrease the wear resistance.

Description

A binder-free cemented carbide and its preparation method and application Technical Field The invention relates to the technical field of cemented carbide, and in particular to a cemented carbide without a binder phase and a preparation method and application thereof. Background Art Traditional WC-Co (Ni, Fe)-other additives cemented carbide has been widely used in many fields such as cutting tools, molds, mining tools, and wear-resistant parts due to its good comprehensive properties such as strength, hardness, toughness and wear resistance. Traditional cemented carbide is mainly prepared by powder metallurgy, that is, WC powder, binder phase powder (mainly including Co, Ni or Fe, whose weight content is generally 6wt.%~25wt.%), and other additive powders are ball milled, mixed, molded and sintered. During the sintering process, fine WC particles (or grains) are dissolved 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 carbide are usually in the shape of triangular prisms. In order to improve the hardness and toughness of traditional cemented carbide, technicians in this field use the anisotropic characteristics of the hardness of the {0001} end crystal plane of HCP structure WC, which is much higher than that of the {1100} prism crystal plane, to transform the morphology of WC grains from the usual triangular prism to a plate (sheet) shape, that is, to promote the lateral growth of the prism plane of WC grains and increase the proportion of the {0001} crystal plane, thereby obtaining a cemented carbide with high hardness and toughness. In order 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 vertical direction, thereby forming plate (sheet) shaped WC grains. However, the above schemes have the following disadvantages: (1) The flattening process of the raw material powder is complicated; (2) During the sintering process, the plate-like grains formed are often thicker (generally greater than 0.8 μm) due to the vertical growth; (3) Due to the dissolution-precipitation growth mechanism, almost all plate-like grains are formed, which is difficult to quantify. Control the content of plate-like grains. In addition, due to the inherent characteristics of high binder phase metal content, the toughness of traditional cemented carbide is better, and the 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 binder phase has been developed, which refers to 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 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 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 pipes, 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 sub-micron 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 it is difficult to meet the application environment with impact loads, especially in the application of cutting tools. Therefore, how to improve the toughness of binderless cemented carbide is an urgent problem to be solved in this field. In view of this, the present invention is proposed. Summary of the invention The purpose 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. 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. The present invention is achieved in that: In a first aspect, the present invention provides a binderless cemented carbide, wherein the cemented carbide is mainly composed of a hard phase; the grain morphology of the hard phase includes ultrafine grains and lamellar grains, and the average grain size of the ultrafine grains is 0.2μm~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 10~60:1. In an optional embodiment, in the binder-free cemented carbide, the volume content of the platelet crystals is 10% to 70%; 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; Preferably, the plate-like crystals are uniformly distributed in the polycrystalline matrix formed by ultrafine crystals. In an optional embodiment, the chemical composition includes, by mass percentage, 0% to 1% bonding 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 remainder is tungsten and unavoidable impurities. In an optional embodiment, the chemical composition includes, by mass percentage, 0.1% to 0.6% of binder phase metal, 0.01% to 0.03% of boron, 1% to 2% of titanium, 1.5% to 2.5% of tantalum, 0.5% to 1.5% of niobium, 0.9% to 1.1% of chromium, 0.2% to 0.4% of vanadium, 6.423% to 6.653% of carbon, and the balance is tungsten and unavoidable impurities; Preferably, the binder phase metal is selected from at least one of cobalt, nickel and iron. 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, mixing the raw materials to obtain a powdery mixture; The powdery mixed material is subjected to a sintering pretreatment to obtain a green body to be sintered, and the green body to be sintered is sintered. In an optional embodiment, the preparation process of the powdered mixture includes: mixing raw material powders weighed according to the chemical composition by wet ball milling, and then drying; 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; Preferably, the binder phase metal powder is a single metal powder or alloy powder. When the powder is cobalt powder, the content of close-packed hexagonal crystal structure in the powder is greater than 90wt.%; 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 to 0.8; Preferably, the tungsten boride powder is WB powder; 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 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. In an optional embodiment, the process of mixing by wet ball milling includes: firstly ball milling the other powders except the binder metal powder and the forming agent powder for 24h to 36h, then adding the binder metal powder and ball milling for 24h to 36h, and then adding the forming agent powder and ball milling for 15h to 20h; Preferably, the molding agent powder is at least one of paraffin wax and polyethylene glycol; Preferably, the ball-to-material mass ratio used in wet ball milling is (4-6):1; Preferably, the medium used in wet ball milling is alcohol. In an optional embodiment, the process of pre-sintering treatment includes: molding the powdered mixture to obtain a primary body, and degreasing the primary body. In an optional embodiment, the process of sintering the green body to be sintered includes: firstly vacuum sintering the green body to be sintered to obtain a pre-sintered body, and then hot isostatic pressing the pre-sintered body; Preferably, the vacuum sintering temperature is 1540°C to 1580°C; the holding time is 1h to 3h, and the pressure is less than or equal to 30Pa; Preferably, the hot isostatic pressing sintering temperature is 1400°C to 1450°C, the holding time is 1h to 3h, and the pressure is 150MPa to 300MPa; More preferably, the hot isostatic pressing sintering is performed under an inert atmosphere. In a third aspect, the present invention provides use of any binderless cemented carbide according to the aforementioned embodiments or a binderless cemented carbide prepared by any of the aforementioned preparation methods in the preparation of cutting tools, molds or mining tools. The present invention has the following beneficial effects: The content of the binder phase metal is less than or equal to 1wt.%, 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 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 part of the grains to form an ultra-thin flake morphology, and achieving the preparation of polygonal flake morphology grains under the condition 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 the hardness and thus reducing its wear resistance. BRIEF DESCRIPTION OF THE DRAWINGS In order to more clearly illustrate the technical solutions of the embodiments of the present invention, the drawings required for use in the embodiments are briefly introduced below. It should be understood that the following drawings only show certain embodiments of the present invention and therefore should not be regarded as limiting the scope. For ordinary technicians in this field, other related drawings can be obtained based on these drawings without creative work. FIG1 is a grain morphology image (low magnification image) of a fracture of a cemented carbide without a binder phase prepared in Example 1.2 provided by the present invention; FIG2 is a grain morphology image (high magnification image) of a fracture of a cemented carbide without a binder phase prepared in Example 1.2 provided by the present invention; FIG3 is a grain morphology diagram of a conventional cemented carbide fracture surface prepared in Comparative Example 1.1 provided by the present invention; FIG. 4 is a grain morphology diagram of the fracture of a binderless cemented carbide prepared in Comparative Example 3.2 provided by the present invention. DETAILED DESCRIPTION In order to make the purpose, technical scheme and advantages of the embodiments of the present invention clearer, the technical scheme in the embodiments of the present invention will be described clearly and completely below. If no specific conditions are specified in the embodiments, the conventional conditions or the conditions recommended by the manufacturer shall be followed. Manufacturers are all conventional products that can be purchased on the market. An embodiment of the present invention provides a binder-free cemented carbide, whose chemical composition, measured by mass percentage, includes: 0% to 1% of binder phase metal, 0.01% to 0.03% of boron, 0.5% to 3% of titanium, 1% to 3% of tantalum, 0% to 2% of niobium, 0.8% to 1.2% of chromium, 0.2% to 0.4% of vanadium, 6.303% to 6.841% of carbon, and the balance is tungsten and unavoidable impurities. It should be noted that the content of the binder phase metal added to the raw material of the binder phase-free cemented carbide provided in the embodiment of the present invention is less than or equal to 1wt.%, and it is difficult for the 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 part of the grains to form a polygonal flake-like morphology. Polygonal flake-like morphology grains can be prepared under the condition of basically no binder phase metal and without the need to flatten the raw material powder, effectively improving the toughness of the binder phase-free cemented carbide without significantly reducing its hardness and thus reducing its wear resistance. 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% 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. In some embodiments, the chemical composition includes, by mass percentage, 0.1% to 0.6% of binder metal, 0.01% to 0.03% of boron, 1% to 2% of titanium, 1.5% to 2.5% of tantalum, 0.5% to 1.5% of niobium, 0.9% to 1.1% of chromium, 0.2% to 0.4% of vanadium, 6.423% to 6.653% of carbon, and the balance is tungsten and unavoidable impurities. The content of each element is optimized to further improve the toughness and hardness of the alloy. 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 the above. The binder-free cemented carbide is mainly composed of a hard phase, and the initial binder phase metal or alloy component reacts with other components in situ to form a hard phase. Here, "mainly" means that it is basically a hard phase, and specifically, the mass percentage of the hard phase reaches more than 99%. 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, or the grain morphology of all hard phases may be composed of ultrafine crystals and lamellar crystals, and the lamellar crystals may be uniformly distributed in the polycrystalline matrix formed by the ultrafine crystals. In the grain morphology of the cemented carbide, the volume content of the lamellar crystals may be 10% to 70%, such as 10%, 20%, 30%, 40%, 50%, 60%, 70%, etc. Due to the presence of the lamellar crystals, the fracture toughness of the alloy is greatly increased. 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. In some embodiments, the flaky crystals are in the form of polygonal thin flakes, the polygonal equivalent circle diameter of the flaky 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 flaky crystals is 10 to 60:1. The polygonal equivalent circle diameter is more than 10 times the thickness, and the flaky crystals are very thin. Specifically, the polygonal equivalent circle diameter of the lamellar 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. The present invention also provides a method for preparing a binderless cemented carbide, comprising: preparing the binderless cemented carbide according to the chemical composition of the binderless cemented carbide, so that the chemical composition of the prepared binderless cemented carbide meets the requirements of the element composition of the binderless cemented carbide. Specifically, the method may include the following steps: S1. Mixing 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 use. 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. In some embodiments, the raw material powder includes: bonding 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. It should be noted that chromium carbide powder and vanadium carbide powder help to inhibit the growth of WC grains by grain boundary diffusion. Compared with traditional cemented carbide, due to the extremely low content of the binder phase, the ultra-hard boride formed further reduces the dissolution of WC in the binder phase metal, and the probability of WC grains growing by the dissolution-precipitation mechanism in the binder phase metal is very low. Most of them are achieved by inhibiting the grain boundary diffusion of WC grains by chromium carbide or vanadium carbide particles, thereby achieving the effect of grain refinement, thereby forming a polycrystalline matrix with an average grain size of 0.2μm to 0.4μm ultrafine grains. At the same time, the interaction between chromium and vanadium elements and boron elements will also affect the morphological characteristics of the lamellar crystals and the uniformity of element distribution. 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. In some embodiments, when the binder phase metal powder is cobalt powder, the content of close-packed hexagonal crystal structure (HCP) in the powder is greater than 90wt.%. 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. 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) surface of the WC crystal, thereby hindering the diffusion growth of the (0001) crystal surface of the WC grain; that is, it can inhibit the WC grain from spreading along the

[0001] In addition, compared with 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 is also beneficial to sintering densification. 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 it easy for WB to react with the binder metal Co, Ni or Fe in situ to form W. 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 is conducive to the formation of lamellar crystals, which may be due to the fact that 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. 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 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, and other raw material powders refer to other types of powders except tungsten carbide powder and binder metal powder. Specifically, the average particle size of tungsten carbide powder can be 0.1 μm, 0.2 μm, 0.3 μm, 0.4 μm, etc., the average particle size of 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. In some embodiments, the process of mixing by wet ball milling includes: firstly, ball milling other powders (i.e., hard phase powders) except the binder phase metal powder and the molding agent powder for 24 to 36 hours, then adding the binder phase metal powder and ball milling for 24 to 36 hours, and then adding the molding agent powder and ball milling for 15 to 20 hours. The raw material powder is added step by step to perform rolling wet ball milling, so that the particle size of the raw material powder is The powder size becomes ultra-fine, which improves the activity of the raw material powder on the one hand and the mixing uniformity of the raw material powder on the other hand. It should be noted that firstly, the powders other than the binder metal powder and the molding agent powder are ball-milled and mixed, so as to avoid the influence of the plastic binder metal powder on the brittle hard phase powder during ball milling, and to improve the crushing efficiency and mixing uniformity, so as to promote the formation of highly active ultrafine particles, which is helpful for the subsequent sintering to control the growth of lamellar crystals, and also helps to form ultrafine grain structure. Finally, the molding agent powder is added and mixed, so as to avoid the influence of the molding agent powder on the ball milling process. Since the molding 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 easy to mix unevenly, but also easily cause the molding agent itself to agglomerate and mix unevenly. Therefore, in the embodiment of the present invention, the hard phase powder is first crushed and mixed evenly, and then mixed evenly with the binder metal powder, and finally mixed with the molding agent powder. 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. 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. 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. S2. Pre-sintering treatment The powdery mixed material is subjected to a pre-sintering treatment to obtain a green body to be sintered, and the specific process of the pre-sintering treatment is not limited. In some embodiments, the pre-sintering treatment process includes: molding the powdered mixed material to obtain a primary body, and degreasing the primary body to remove the molding agent in the body. S3, Sintering The green body to be sintered can be sintered by combining vacuum sintering and hot isostatic pressing. The formula is beneficial to improve the toughness and hardness of the prepared cemented carbide without a binder phase. In the actual operation process, the process of sintering the green body to be sintered includes: firstly 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 cemented carbide without a binder phase. In some embodiments, the vacuum sintering temperature is 1540℃~1580℃; the holding time is 1h~3h, and the pressure is less than or equal to 30Pa; by selecting a suitable vacuum sintering temperature, it is helpful to promote sintering densification without causing abnormal grain growth. If the vacuum sintering temperature is lower than 1540℃, the pre-sintered body contains more open pores, which cannot meet the density requirements of the pre-sintered body in the next step of hot isostatic pressing sintering; if the vacuum sintering temperature is higher than 1580℃, 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 cemented carbide without a binder phase. 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. In some embodiments, the hot isostatic pressing sintering temperature is 1400°C to 1450°C, the holding time is 1h to 3h, the pressure is 150MPa to 300MPa, and the hot isostatic pressing sintering is performed under an inert atmosphere (such as argon). By optimizing the hot isostatic pressing sintering parameters, the residual pores in the pre-sintered body are eliminated under the premise that the grains in the pre-sintered body no longer grow significantly, and a dense cemented carbide without a binder phase is obtained. 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. It should be noted that the binderless cemented carbide and the preparation method thereof provided in the embodiment of the present invention can prepare a binderless cemented carbide containing ultrafine crystals and lamellar crystals, wherein the lamellar crystals are uniformly distributed in a polycrystalline matrix composed of ultrafine crystals. The fracture toughness is increased. The preparation method provided in the embodiment of the present invention does not need to flatten the raw material powder during the preparation process, nor does it need to grow through the dissolution-precipitation mechanism of WC in the bonding phase metal, and can form lamellar crystals. The preparation method is simple and effective. In addition, the non-bonding phase cemented carbide prepared in the embodiment of the present invention retains an ultrafine-grained polycrystalline matrix, 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 non-bonding phase cemented carbide with high toughness and hardness is obtained. 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, and utilize its high toughness and hardness to make the prepared cutting tools, molds, mining tools and other products also have excellent performance. The organizational structure characteristics and properties of the binderless cemented carbide prepared by the present invention are further described in detail below in conjunction with the examples. Experimental Group 1 Test group 1 includes Examples 1.1 to 1.5 and Comparative Examples 1.1 to 1.2, and mainly adjusts the content of Co, as follows: Example 1.1 This embodiment provides a method for preparing a binder-free cemented carbide, which is prepared by the following steps: (1) The raw material powders are accurately weighed according to the proportion, and mixed by rolling wet ball milling with stepwise addition of the raw material powders, with a ball-to-material ratio of 5:1, and the medium used in the wet ball milling is alcohol; and the mixed material powder is prepared by spray drying. Among them, the composition of the raw material powders, by mass percentage, is as follows: (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. 0.75 W 0.25 The morphology of raw material powders of )C, TaC, NbC, WB, Cr3C2 and VC are all equiaxed. The process of adding raw material powders step by step includes: firstly, ball milling other brittle powders except the binder metal cobalt powder and the molding agent for 30 hours; then adding the binder metal cobalt powder for 30 hours (when the initial binder metal powder content is 0, this step is omitted); finally, adding the paraffin molding agent powder and ball milling for 18 hours. (2) The blank is pressed by molding. (3) Degreasing the obtained green body is performed using an argon negative pressure process. (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. (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. Example 1.2-Example 1.5 The only difference between Example 1.2 to Example 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. Comparative Example 1.1-Comparative Example 1.2 The difference between Comparative Examples 1.1 and 1.2 and Example 1 is that the composition of the raw material powder and the sintering process of the cemented carbide are adjusted. The composition of the raw material powder in the cemented carbide is shown in Table 1. The sintering process used in Comparative Examples 1.1 and 1.2 is as follows: The vacuum sintering temperatures used in Comparative Examples 1.1 and 1.2 are 1420°C and 1400°C respectively, the holding time is 2h, and the pressure is less than or equal to 30Pa; then a low-pressure sintering process is adopted, and the sintering temperatures of Comparative Examples 1.1 and 1.2 are 1400°C and 1380°C respectively, the holding time is 2h, and the argon pressure is 5MPa. Note: Comparative Examples 1.1 and 1.2 are the formula compositions of conventional cemented carbides with a high cobalt content. If the same vacuum sintering temperature (eg, 1560°C) as that of cemented carbides without a binder phase is used, severe overburning will result. Table 1 Chemical formula and ratio of raw material powder used in test group 1 (wt.%) The content of each chemical component in the prepared cemented carbide was tested, and the test methods used were as follows: the content of metal elements was tested with reference to the national standard GB / T26050-2010 "Determination of metal element content in cemented carbide by X-ray red fluorescence melting method"; the total carbon content was tested with reference to the national standard GB / T42275-2022 "Determination of total carbon content in cemented carbide by high-frequency combustion infrared absorption method / thermal conductivity method"; the boron content was determined with reference to the national standard GB / T39138.3-2020 "Chemical analysis method of gold-nickel-chromium-iron-silicon-boron alloy Part 3: Determination of chromium, iron, silicon and boron content by inductively coupled plasma emission spectrometry". The test results are shown in Table 2. Table 2 Chemical composition ratio of cemented carbide in test group 1 (wt.%) The microstructure and properties of the cemented carbides obtained by testing Example 1.1 to Example 1.5 and Comparative Example 1.1 to Comparative Example 1.2 are shown in Table 3. Test method: (1) Test of the morphology and size of the flake crystals: The test was conducted by SEM observation, as follows: The morphology characteristics were observed by SEM; at least 500 flake crystals were randomly tested for each sample to obtain the thickness and polygonal area of ​​each flake crystal. The polygons were equivalent to circles using the equal area method, and the equivalent circle diameter was obtained using the following formula: D i =(4S i / π) 1 / 2 In the formula, S i is the measured value of the polygonal area of ​​the ith platelet; D i is the polygonal equivalent circle diameter of the i-th lamellar crystal. (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 test volume, the volume content of the flake crystals is calculated. (3) Test of ultrafine grain size of matrix: Tested by SEM observation. The test method refers to the national standard GB / T3488.2-2018 "Metallographic determination of cemented carbide microstructure Part 2: Measurement of WC grain size". (4) The tests of fracture toughness and hardness refer to 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" respectively. Table 3 Comparison of the organizational structure and performance parameters of cemented carbide obtained in test group 1 It can be seen from Table 3 that the fracture toughness of the binder-free 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 of the binder phase increases, the atomic diffusion capacity increases, the polygonal equivalent circle diameter of the formed flake grains increases, the thickness also increases, the ratio of the polygonal equivalent circle diameter to the thickness decreases, and the content of the flake grains increases. The amount of MgO also increases slightly; correspondingly, the size of the ultrafine grains in the matrix also increases slightly. The typical organizational structure is shown in Figures 1 and 2. Although the fracture toughness of the cemented carbides obtained in Comparative Examples 1.1 and 1.2 is also higher than 7 MPa·m 1 / 2 , but the hardness is only below 2108. The Co content of Comparative Examples 1.1 and 1.2 is high. Due to the enhanced dissolution-precipitation of WC in the liquid phase cobalt during sintering, in the absence of flattened raw material powder particles as crystal nuclei guide, even with the presence of boron, conventional triangular prism-shaped grains are still formed, and no lamellar crystals are formed. Figure 3 shows the grain morphology of the cemented carbide fracture obtained in Comparative Example 1.1. On the one hand, due to the increase in the cobalt content of the bonding phase, the fracture toughness of the cemented carbide increases and the hardness decreases; on the other hand, with the increase in the cobalt content of the bonding phase, the morphology of the lamellar crystal changes from lamellar to columnar, which will lead to a decrease in the fracture toughness of the alloy; therefore, the comprehensive results show that when the cobalt content of the bonding phase reaches 6%, compared with Example 1.5, the toughness of Comparative Examples 1.1 and 1.2 is not significantly increased, but the hardness decreases significantly, with the hardness decreasing by 9% to 16%. Experimental Group 2 Experimental Group 2 includes Example 1.2, Example 2.1-Example 2.2, and Comparative Example 2.1-Comparative Example 2.2, mainly adjusting the content of boron element, and correspondingly adjusting the type and ratio of raw material powders, as shown in Table 4: Table 4 Chemical formula and ratio of raw material powder used in test group 2 (wt.%) The preparation methods of Examples 2.1 to 2.2 and Comparative Examples 2.1 to 2.2 refer to those of Example 1.1. 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. Table 5 Chemical composition ratio of cemented carbide without binder phase in test group 2 (wt.%) The microstructure and properties of the binder-free cemented carbide were obtained by testing Example 1.2, Example 2.1-Example 2.2, and Comparative Example 2.1-Comparative Example 2.2, as shown in Table 6. Table 6 Comparison of the microstructure and performance parameters of the binder-free cemented carbide obtained in test group 2 It can be seen from Table 6 that when no boron is added, no lamellar grains are formed, and an ultrafine-grained cemented carbide structure without a binder phase is obtained, with high hardness but low toughness. When boron is added, the hardness decreases due to the formation of lamellar crystals, but the toughness increases. With the increase of boron content, on the one hand, the ratio of the polygonal equivalent circle diameter to thickness of the lamellar crystal increases significantly, the content of the lamellar crystal also increases, and its toughness increases; on the other hand, due to the formation of W a Co b B c The superhard compound phase improves the disadvantage of low hardness of the bonding phase metal cobalt, and its hardness is also increased. However, if the amount of boron added is too high, the grains are prone to grow, that is, the size of the lamellar crystals and the size of the matrix grains increase, which will lead to a decrease in the fracture toughness and hardness of the alloy. Experimental Group 3 Experimental group 3 includes Example 1.2, Example 3.1-Example 3.4, and Comparative Examples 3.1-Comparative Examples 3.3, mainly adjusting the content of titanium element, and correspondingly adjusting the type and ratio of raw material powders, as shown in Table 7: Table 7 Chemical formula and ratio of raw material powder used in test group 3 (wt.%) The preparation methods of Examples 3.1 to 3.4 and Comparative Examples 3.1 to 3.3 refer to Example 1. The chemical composition test results of the binder-free cemented carbides of Example 1.2, Example 3.1 to Example 3.4, and Comparative Example 3.1 to Comparative Example 3.3 are shown in Table 8. Table 8 Chemical composition ratio of test group 3 cemented carbide without binder phase (wt.%) Test Example 1.2, Example 3.1-Example 3.4, and Comparative Example 3.1-Comparative Example 3.3 to obtain The microstructure and properties of cemented carbide without binder phase are shown in Table 9. Table 9 Comparison of the microstructure and performance parameters of the binder-free cemented carbide obtained in test group 3 From Table 9, we can see that with (Ti 0.75 W 0.25 )C content 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 is a grain morphology of the fracture of comparative example 3.2. When the content of plate crystals is about 40vol.% to 60vol.%, and the average particle size of the matrix ultrafine crystals is 0.2μm to 0.4μm, the comprehensive performance of the alloy reaches the best. With the increase of (Ti 0.75 W0.25 ) As the C content continues to increase, on the one hand, the content of lamellar crystals will exceed 60 vol.%, and the fracture toughness of the alloy will decrease due to the bridge phenomenon between the lamellar crystals. 0.75 W 0.25 ) When the addition amount of C is too high, the grains are easy to grow, 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. Experimental Group 4 Experimental group 4 includes Example 1.2, Example 4.1-Example 4.4 and Comparative Example 4.1-Comparative Example 4.3, mainly adjusting the content of tantalum or niobium element, and correspondingly adjusting the type and ratio of raw material powder, as shown in Table 10: Table 10 Chemical formula and ratio of raw material powder used in test group 4 (wt.%) The preparation methods of Examples 4.1 to 4.4 and Comparative Examples 4.1 to 4.3 refer to Example 1. The chemical composition test results of the binder-free cemented carbides obtained from Example 1.2, Example 4.1 to Example 4.4, and Comparative Example 4.1 to Comparative Example 4.3 are shown in Table 11. Table 11 Chemical composition ratio of test group 4 cemented carbide without binder phase (wt.%) The microstructure and properties of the cemented carbide 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. Table 12 Comparison of the microstructure and performance parameters of the binder-free cemented carbide obtained in test group 4 As shown in Table 12, the change in the content of TaC and NbC has little effect on the formation of lamellar crystals, but as the content of TaC and NbC decreases, the content of WC needs to be increased accordingly, which causes the hardness of the alloy to increase and the fracture toughness to decrease slightly. The main purpose of adding TaC and NbC is to help improve the oxidation resistance of the alloy. Experimental Group 5 Test group 5 includes Example 1.2, Example 5.1-Example 5.4, and Comparative Example 5.1-Comparative Example 5.4, mainly adjusting the content of chromium or vanadium elements, and correspondingly adjusting the types and proportions of raw material powders, as shown in Table 13: Table 13 Chemical formula and ratio of raw material powder used in test group 5 (wt.%) The preparation methods of Examples 5.1 to 5.4 and Comparative Examples 5.1 to 5.4 refer to Example 1. The chemical composition test results of the binder-free cemented carbides obtained from Example 1.2, Example 5.1 to Example 5.4, and Comparative Example 5.1 to Comparative Example 5.4 are shown in Table 14. Table 14 Chemical composition ratio of test group 5 cemented carbide without binder phase (wt.%) The microstructure and performance of 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. Table 15 Comparison of the microstructure and performance parameters of the binder-free cemented carbide obtained in test group 5 It can be seen from Table 15 that when the content of Cr3C2 and VC is low, the growth inhibition of the matrix grains is weak, and the matrix grains grow significantly, resulting in a lower hardness of the alloy; with the increase of Cr3C2 and VC content, the size of the lamellar crystals decreases, the ratio of the polygonal equivalent circle diameter / thickness of the lamellar crystals decreases, and the size of the matrix grains also decreases, while the content of the lamellar crystals has little effect. Therefore, the hardness of the alloy increases and the toughness decreases; but when the content of Cr3C2 and VC is too high, agglomeration is easy to occur during the sintering process, which may be caused by the presence of boron elements that easily attract Cr elements to agglomerate, thereby increasing the brittleness of the alloy, and the fracture toughness shows a trend of first increasing and then decreasing. Experimental Group 6 This comparative example 6.1 provides a binder-free cemented carbide and a preparation method thereof, which differs from Example 1.2 only in that: 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. The microstructure and properties of the binder-free cemented carbide obtained in Comparative Example 6.1 and Example 1.2 are shown in Table 16. Table 16 Comparison of properties and structural parameters of cemented carbide without binder phase It can be seen from Table 16 that the direct use of TiC powder will cause the temperature required for sintering to become dense to increase. Due to the low vacuum sintering temperature in Example 6.1, most of the residual pores in the pre-sintered body are 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 performed. Test Group 7 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 raised to The temperature is 1600°C (Comparative Example 7.2), and the rest is the same as Example 1.2. 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. Table 17 Comparison of properties and structural parameters of cemented carbide without binder phase As can be seen from Table 17, when the vacuum sintering temperature is lowered, the residual pores in the pre-sintered body of Comparative Example 7.1 are mostly open pores due to the low vacuum sintering temperature, and the density requirement of the pre-sintered body in the next step of hot isostatic pressing is not met, so the material is not sintered densely, and further testing is not performed. When the vacuum sintering temperature is increased, the grains grow significantly, and the ratio of the polygonal equivalent circle diameter / thickness of the plate-like crystals in Comparative Example 7.2 decreases sharply, and the matrix grains have grown to the micron level, which ultimately leads to a sharp decrease in fracture toughness and hardness. The above are only preferred embodiments of the present invention and are not intended to limit the present invention. For those skilled in the art, the present invention may have various modifications and variations. Any modification, equivalent replacement, improvement, etc. made within the spirit and principle of the present invention shall be included in the protection scope 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 grain size of the ultrafine crystals is 0.2 μm to 0.4 μm; the lamellar crystals are in the shape of polygonal flakes, and the ratio of the polygonal equivalent circle diameter to the thickness of the lamellar crystals is 10 to 60:

1.

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%; 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; Preferably, the lamellar crystals are uniformly distributed in the polycrystalline matrix formed by the ultrafine crystals.

3. The binder-free cemented carbide according to claim 1, characterized in that: Calculated by mass percentage, its chemical composition includes: bonding 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.

4. The binder-free cemented carbide according to claim 3, characterized in that: Its chemical composition, in terms of mass percentage, includes: 0.1% to 0.6% of binder metal, 0.01% to 0.03% of boron, 1% to 2% of titanium, 1.5% to 2.5% of tantalum, 0.5% to 1.5% of niobium, 0.9% to 1.1% of chromium, 0.2% to 0.4% of vanadium, 6.423% to 6.653% of carbon, and the balance is tungsten and unavoidable impurities; Preferably, the binder phase metal is selected from at least one of cobalt, nickel and iron.

5. A method for preparing a binder-free cemented carbide according to any one of claims 1 to 4, characterized in that: include: The ingredients are prepared according to the chemical composition of the binderless cemented carbide, and the raw materials are mixed to obtain a powdery mixture; The powdery mixed material is subjected to a sintering pretreatment to obtain a green body to be sintered, and the green body to be sintered is sintered.

6. The preparation method according to claim 5, characterized in that: The preparation process of the powdered mixed material comprises: mixing raw material powders weighed according to chemical composition by wet ball milling, and then drying; 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; Preferably, the binder phase metal powder is a single metal powder or an 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 90wt.%; 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 to 0.8; Preferably, the tungsten boride powder is WB powder; Preferably, the overall average particle size of the raw material powder 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 phase 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.

7. The preparation method according to claim 6, characterized in that: The process of mixing by wet ball milling includes: firstly ball milling 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; Preferably, the molding agent powder is at least one of paraffin and polyethylene glycol; Preferably, the ball-to-material mass ratio used in the wet ball milling is (4-6):1; Preferably, the medium used in the wet ball milling is alcohol.

8. The preparation method according to claim 5, characterized in that: The process of the pre-sintering treatment includes: molding the powdery mixed material to obtain a primary body, and degreasing the primary body.

9. The preparation method according to claim 5, 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; Preferably, the vacuum sintering temperature is 1540°C to 1580°C; the holding time is 1h to 3h, and the pressure is less than or equal to 30Pa; 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; More preferably, the hot isostatic pressing sintering is performed under an inert atmosphere.

10. Use of the binderless cemented carbide according to any one of claims 1 to 4 or the binderless cemented carbide prepared by the preparation method according to any one of claims 5 to 9 in the preparation of cutting tools, molds, or mining tools.

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