JPS635350B2 - - Google Patents

Description

[Detailed description of the invention]

The present invention relates to a dense ceramic sintered body made of refractory hard materials such as carbides, borides, and nitrides of metals from groups IVa, Va, and VIa of the periodic table, and whose main purpose is to be used as cutting tools. be. Currently used sintered bodies for cutting tools that are mainly made of refractory hard materials include cemented carbide in which the main material, tungsten carbide, is bonded with cobalt metal, and titanium carbide, the main material, in which nickel metal is bonded. cermet, alumina-based ceramic, and cubic boron nitride (hereinafter referred to as CBN), which is the main material.
) or a sintered body made by adding a binder to diamond and densifying it under ultra-high pressure. The dense sintered body of the present invention is characterized in that it is composed of two or more types of refractory hard materials and contains substantially no binding metal. Therefore, among the above sintered bodies,
Since cemented carbide and cermet are alloys in which the main material is a refractory hard material bonded with metal, their constituent phases are essentially different from the sintered body of the present invention. The above-mentioned sintered body for a cutting tool is required to be a completely or almost completely dense sintered body from the viewpoint of required edge strength. Regarding the manufacturing method, first, since cemented carbide and cermet are liquid-phase sintered alloys, the green compact is easily densified by the surface tension of the liquid phase generated during the sintering process. Therefore, dense sintered bodies of these alloys can be produced at low cost by ordinary sintering at 1350°C to 1500°C. On the other hand, densification of ceramics based on solid phase sintering is more difficult than that of liquid phase sintered alloys.
Since solid phase sintering is fundamentally based on diffusion phenomena, the temperature required for densification in solid phase sintering depends primarily on the melting temperature of the material. Therefore, solid phase sintering of high melting point materials with melting temperatures of 3000°C or higher, such as carbides such as titanium, zircon, hafnium, tantalum, and niobium, requires high temperatures of 2000°C or higher. In addition to equipment and technical difficulties, grain growth associated with high-temperature sintering deteriorates the properties of the sintered body. Therefore, to date, it has been considered impossible to produce dense ceramics based on high-melting-point substances by ordinary sintering methods, and of course they have not been put to practical use as materials for cutting tools. As described above, since it is generally difficult to densify a refractory hard material by solid-phase sintering, various methods have been used to produce a dense sintered body. One example is a method of promoting densification by adding a small amount of a low melting point substance or a substance that forms a low melting point substance during the sintering process. For example, sintering alumina often involves adding small amounts of foreign oxides such as MgO. However, these additives or substances formed by these additives remaining in the sintered body after sintering generally significantly deteriorate the properties of the sintered body. Therefore, hot presses are more widely used in the production of dense sintered bodies of refractory hard materials. For example, a sintered body made of 70% by weight alumina and 30% by weight titanium carbide is an alumina-based ceramic for cutting tools whose use has recently been expanded, but a hot press is essential for its production. The hot pressing method is significantly inferior in productivity to the normal sintering method, and the manufacturing cost is high due to wear and tear of dies and the like.
Still another method is to completely densify the material by hot isostatic pressing after normal sintering. Sintered bodies based on diamond or CBN generally contain high melting point metals such as tungsten and molybdenum, titanium carbide, titanium nitride,
Solid phase sintering is performed using a refractory hard material such as alumina as a binder. In this case, to prevent transformation of diamond or CBN during the sintering process,
It must be sintered under ultra-high pressure of over 45,000 atmospheres. Despite using a high-melting point binder, the green compact is heated to temperatures between 1200℃ and 1400℃ due to the action of this ultra-high pressure.
Easily densified at relatively low temperatures. However, since these sintered bodies require ultra-high pressure equipment for the production and sintering of the main material, they are much more expensive than other sintered bodies. As mentioned above, ceramics for cutting tools are currently manufactured using methods that are generally expensive to manufacture, but the present invention provides a dense ceramic for cutting tools made of refractory hard materials such as carbides, nitrides, and borides. This invention relates to a ceramic sintered body manufactured at low cost by a common sintering method. Next, the cutting performance of the sintered body for cutting tools will be explained. Typical cutting performance characteristics that a cutting tool material should have are wear resistance and chipping resistance. However, with the recent remarkable progress in processing technology, the demand for higher efficiency in cutting processes has increased, and cutting tool materials are being forced to cope with unprecedentedly severe cutting conditions. Typical examples of this are increasing cutting speeds and making work materials more difficult to cut, especially harder, but it is not possible to adequately meet these demands by simply improving the cutting performance characteristics of conventional cutting tool materials. Therefore, there is a strong demand for the development of new tool materials. The practical application of expensive CBN sintered bodies is an example of meeting this requirement. To explain the situation specifically, first of all, it is difficult to increase the cutting speed even though cemented carbide and cermet containing bonding metals have the same capability. For example, when cutting carbon steel, the upper limit of the general practical cutting speed for cemented carbide is 150 m/min, and for cermet it is about 250 m/min. Regarding these liquid phase sintered alloys, improvements in the heat resistance properties of the bonding metal phase are being considered as a countermeasure for increasing speed, but it is still difficult to completely overcome the inherent defects of the bonding metal in order to increase speed. That is impossible, and after all
It is thought that 300m/min is the limit of practical cutting speed. Ceramics, on the other hand, are extremely promising for responding to higher speeds. Even with current alumina-based ceramics, it is possible to put this into practical use at cutting speeds of 400 m/min to 500 m/min. However, CBN sintered bodies do not necessarily show excellent performance in high-speed cutting of soft materials. On the other hand, cermets are not as good at responding to hardening of the work material due to the problem of plastic deformation of the cutting edge. Alumina-based ceramics have excellent performance in terms of welding resistance with work materials, high-temperature strength, and high-temperature hardness, but are associated with low thermal conductivity at high temperatures and have the problem of cracking due to thermal shock. Therefore, the hardness of the workpiece material is limited to about HRC55, making it difficult to put it to practical use in cutting hardened steel with HRC60 or higher. Therefore, the cutting speed of these hardened steels is conventionally 20m/min to 30m/min using cemented carbide.
Low speed cutting at n was the only method. In response to this situation, the advent of CBN sintered bodies that can cut hardened steel at speeds of 50 m/min to 100 m/min has completely changed the concept of cutting hard materials, and will improve many fields of grinding in the future. The possibility of replacing it with efficient cutting is also emerging. In this case, the extremely high price of CBN sintered bodies would be the biggest practical problem. As is clear from the above explanation, specifically, the cutting performance in high-speed cutting and cutting of high-hardness materials is superior to that of conventional alumina ceramics, and preferably the cutting performance of the latter is similar to that of CBN sintered bodies.
There is a strong demand for the emergence of inexpensive cutting tool materials. The present invention relates to an inexpensive, high-performance ceramic for cutting tools and a method for manufacturing the same, aimed at meeting this demand. By the way, in order to cope with higher cutting speeds or harder workpiece materials, tool materials need to have better high-temperature properties than conventional materials. High-temperature properties that a cutting tool material should possess include (1) hardness, (2) strength and toughness, (3) chemical stability such as welding resistance with the work material and oxidation resistance, and (4)
It has thermal shock resistance. Refractory hard substances such as carbides, borides, and nitrides of transition metals in groups IVa, Va, and VIa of the periodic table have long been attracting attention as high-temperature materials, especially cutting tool materials, and are used as main materials for cemented carbide and cermets. Although these substances have been put into practical use, considering the high-temperature properties mentioned above, these substances cannot be used as main materials for cutting tool materials.
Even today, it remains an extremely promising substance. The biggest reason is that these materials generally have high room temperature hardness and high melting point,
This is due to the fact that it has a metallic bond. That is, the former ensures high hardness and strength at high temperatures, and the latter guarantees superior thermal shock resistance compared to other ceramics such as oxides due to its good thermal conductivity. The excellent high-temperature strength of these refractory hard materials is based on the following grounds. In other words, the strength of a material is determined by the stress that causes it to break, whereas the hardness of a material depends on how easily it deforms in response to stress. As is clear from the fact that both of these properties are expressed in units of stress, there is a certain proportional relationship between the two properties in substances with the same type and form of chemical bonding. In brittle materials, localized fracture accompanied by cracking occurs during hardness measurements, and it has been experimentally confirmed that almost the same relationship exists between strength and hardness in this case as well. On the other hand, the recrystallization phenomenon of polycrystals means an increase in the mobility of atoms, and recrystallization reduces deformation resistance. Moreover, this recrystallization temperature is related to the melting point of the material. Therefore,
Materials with high cold hardness and high melting point can be regarded as materials with inherently good high temperature hardness and high temperature strength.

【table】

[Table] Table 1 shows the crystal form, melting point, and Vickers hardness of the main carbides, borides, and nitrides of group IVa, Va, and VIa metals. The numerical values are mainly based on the literature of Samsonov et al. (GVSamsonov and IM
Vinitskii: Handbook of Refractory
Compounds, IFI/PLENUM NEW YORK,
(1980). However, nitrides of Group VIa metals have rather low melting points and Vickers hardnesses, and therefore do not belong to the category of refractory hard materials.
Therefore, since it is outside the scope of the present invention, it has been omitted. When comparing these refractory hard materials from the viewpoint of high-temperature strength as mentioned above, it is found that all compounds of carbides, borides, and nitrides have group IVa and Va
It is pointed out that while compounds of group metals generally have excellent properties, compounds of group VIa metals, especially chromium compounds, have relatively low hardness and melting point. For the above reasons, carbides, borides, and nitrides of IVa and Va group metals and carbides and borides of group VIa metals (hereinafter referred to as the refractory hard materials), excluding chromium compounds, generally have excellent high-temperature strength. Be expected. However, these materials have another problem, which is a decrease in strength due to creep. In other words, materials with strong covalent bonds such as Si ₃ N ₄ and SiC are essentially resistant to creep, so the drop in high-temperature strength due to creep is generally expected to be small. For such fire-resistant hard materials that have strong properties, consideration must be given to the decrease in strength due to creep. Creep is a very complex phenomenon, and many mechanisms proceed simultaneously, but it can be broadly classified into plastic deformation at grain boundaries and within grains. Strengthening with solid solutions and heterophase dispersion reinforcement are methods for improving this creep property, and the refractory hard material is extremely advantageous in this respect as well, as described below. As shown in Table 1, all of the IVa and Va group carbides and nitrides have cubic crystal forms, and all borides have hexagonal crystal forms except for chromium boride and molybdenum boride. It is known that when they are isomorphic, these substances are extremely easy to form a solid solution with each other even in different compounds, such as a combination of a carbide and a nitride, and generally form a solid solution. in cermet
The formation of a TiC—TiN solid solution phase is one example.
Although solid solubility is naturally limited between substances of different types, mutual solubility generally exists. for example
Cubic carbides such as WC and Mo ₂ C are solid solution carbides with hexagonal carbides, and WC-
TiC-TaC-NbC solid solution carbide and TiC-Mo ₂ C solid solution carbide are used in cermets.
A further advantage is that the refractory hard materials shown in Table 1 generally tend to coexist relatively stably, even though they have different compounds and crystal forms. Due to these characteristics, when using the refractory hard substance as the main material of a cutting tool material, two or more types of the refractory hard substance must be made into a solid solution during the sintering process, or a solid solution powder must be prepared in advance. It is extremely easy to make the main hard phase into a solid solution by using a method such as using the powder as a raw material powder. On the other hand, by intentionally coexisting two or more hard phases,
It is also possible to expect a structural effect on the properties of the sintered body. Therefore, in ceramics whose main materials are two or more kinds of the refractory hard materials, it is possible to change not only the creep properties but also other properties of the sintered body over a very wide range by the above method. It is possible. What is particularly pointed out here is that among the refractory hard materials shown in Table 1, chromium compounds are unique in terms of crystal form, and Cr ₃ C ₂ , Cr ₇ C ₃ ,
Only Cr ₃ B ₄ and CrB are orthorhombic. This fact suggests that the solid solubility of these orthorhombic chromium compounds with the refractory hard material is probably all limited, but the mutual solubility of these chromium compounds with the refractory hard material is likely to be limited. Regarding TiC
-There has been almost no research on anything other than the Cr ₃ C ₂ system. As detailed above, the fire-resistant hard material is
High-temperature hardness, high-temperature hardness, and
It is still considered to be a very promising material today in terms of high temperature properties such as high temperature strength and thermal shock resistance. This means that as a bonded phase
This is supported by the fact that CBN sintered bodies containing 40% by volume or more of the refractory hard material such as TiN exhibit extremely excellent performance in cutting high-hardness materials. The first idea of the present invention is based on the supposition of the fire-resistant hard material as the main material based on the above-mentioned basis. Despite having such excellent properties, the main reason why ceramics based on the fire-resistant hard material have not been put to practical use as cutting tool materials to date is due to their high melting point, as mentioned above. The problem lies in the difficulty of densification. Therefore, in order to put ceramics mainly made of the refractory hard material into practical use, it is necessary to develop appropriate means for promoting densification. However, for example, in ceramics whose main materials are Si ₃ N ₄ and SiC, which have strong covalent bonds and are difficult to sinter, methods of promoting densification during sintering have been intensively researched and have produced results. On the other hand, the reality is that almost no essential studies have been conducted on promoting the densification of the fire-resistant hard material. Therefore, as a means for promoting densification in the sintering process of the refractory hard material, chromium carbide or chromium boride will be defined as a ceramic binder phase. First, these materials have lower hardness and melting point than the refractory hard materials, and also have considerably larger coefficients of thermal expansion. Furthermore, Cr ₃ C ₂ ,
Orthorhombic compounds such _{as Cr7C3} , _Cr3B4 and _CrB are
Since the refractory hard material and the crystals are similar in shape, solid solution is limited, and as a result, it is expected that the crystals will probably exist as an independent phase even after sintering. These characteristics of chromium compounds are related to the relationship between the main material and the binder phase in other sintered bodies, such as the relationship between the main material and the binder metal in cemented carbide and cermets, or
This is similar to the relationship between CBN and a binder phase such as TiN in a CBN sintered body. From this point of view, chromium compounds, especially orthorhombic chromium carbides and orthorhombic chromium borides, are understood to have the basic properties as a binder phase for the refractory hard material. Next, conventional research on these orthorhombic chromium compounds is extremely limited and sufficient knowledge has not been obtained, but regarding Cr ₃ C ₂ ,
P. Schwarzkopf and R.
kieffer: Refractoky Hard Metals, Mac
Millan., 1953) reveals some interesting facts as follows. The first is
This is the effect of Cr ₃ C ₂ as a diffusion promoter in the preparation of solid solutions such as WC-TiC solid solution carbides. Since solid-phase sintering is essentially based on diffusion phenomena, this effect can be achieved by combining the refractory hard materials, especially two or more such refractory hard materials and Cr ₃ C ₂ as a binder phase.
Assuming a green compact of mixed powder with
It is expected that the action of Cr ₃ C ₂ will be effective in promoting densification. The second is Cr ₃ C ₂ to the refractory hard material.
Another problem is that the addition of metallic chromium has a very remarkable effect on improving the oxidation resistance of the refractory hard material. Among the high-temperature properties that the refractory hard material should possess as a cutting tool material, it has a general drawback of being inherently inferior to oxide-based ceramics in terms of oxidation resistance. To overcome this drawback, the use of Cr ₃ C ₂ as a binder phase improves the oxidation resistance of the sintered body mainly made of the refractory hard material due to the above effects of Cr ₃ C ₂ or metallic chromium. It is expected that this will work extremely effectively. The third and most interesting fact is the research result regarding solid solution of Cr ₃ C ₂ in cubic carbide. As mentioned above, the only previous research on the solid solution of Cr ₃ C ₂ with the refractory hard material is the TiC-Cr ₃ C ₂ system, but Cr ₃ C ₂ is dissolved in TiC in a fairly large amount. for example
It is known that at 1700°C to 1800°C, a solid solution of 40% by weight or more occurs. On the other hand, research on the chromium-carbon system phase diagram shows that CrC clearly exists at high temperatures of 2000°C or higher, but cannot exist at room temperature because it decomposes into Cr ₃ C ₂ and graphite during the cooling process. However, it has been suggested that CrC may exist even at room temperature in NaCl-type face-centered cubic carbides, such as solid solutions in TiC. This suggests the possibility that the orthorhombic chromium compound is solid-dissolved in the same form as the crystal form of the solvent substance when it is dissolved in a similar variant substance whose crystal form is simpler and more stable. Therefore, in the process of sintering a green compact made of a mixture of TiC and Cr ₃ C ₂ , Cr ₃ C ₂ is converted into TiC by CrC.
Assuming that carbon is dissolved in the form of solid solution, the proportion of carbon in the remaining undissolved Cr ₃ C ₂ phase gradually decreases as the solid solution progresses. As a result, it is probably not solidly dissolved.
Near the surface of the Cr ₃ C ₂ phase, it is a lower carbide.
_{Cr7C3 or Cr23C6} will be formed _. Even when Cr ₇ C ₃ is used as the binder phase, it is expected that a reaction similar to that described above will occur since the crystal form is the same as that of Cr ₃ C ₂ . This mechanism also shows that a similar solid solution reaction occurs when the refractory hard substance diboride is used as the main material and orthorhombic chromium boride Cr ₃ B ₄ or CrB is used as the binder phase. expected to do so. i.e. in this case Cr ₃ B ₄ or CrB
Hexagonal chromium boride with the same crystal form as the main material
CrB ₂ forms a solid solution in the main material diboride, and as a result, near the surface of the undissolved Cr ₃ B ₄ phase or CrB phase, lower-order compounds Cr ₅ B ₃ or Cr ₂ B will be formed. The above reaction mechanism in the sintering process and the formation of lower-order compounds near the surface of the undissolved binder phase based on it are not only effective in promoting densification of the green compact based on diffusion solid solution, but also From the viewpoint of the interfacial strength between the main phase and the binder phase in the sintered body,
It is thought that it will play an extremely effective role.
From this point of view, among chromium compounds, orthorhombic chromium carbide and orthorhombic chromium boride (hereinafter referred to as the chromium compounds) are particularly useful as binder phases in ceramics whose main material is the refractory hard substance. It is considered suitable. Regarding the formation of lower-order carbides during the sintering process of orthorhombic carbide Cr ₃ C ₂ , there are research results by Fukatsu (Powder and Powder Metallurgy, VO18.1961.P248) based on a different mechanism. According to Fukatsu, _{Cr3C2 itself}
It is assumed that a portion of Cr ₇ C ₃ and metallic chromium dissociate at temperatures above 800°C, and in Cr ₃ C ₂ -Ni alloys, this dissociation product, Cr ₇ C ₃ and metallic chromium, preferentially converts into the nickel phase. As a result of forming a Ni-Cr alloy through solid solution, it has been confirmed that the Cr ₇ C _three phase does not exist in the sintered body. On the other hand, to prepare a mixed powder of the refractory hard material and the chromium compound as a binder phase, wet mixing is usually performed using a ball mill or a vibration mill. Small amounts of iron group metals are unavoidably present. Therefore, based on the above-mentioned solid solution mechanism as CrC, Cr ₃ C ₂ in the green compact of the mixed powder for sintered body containing this iron group metal,
Even if lower carbides are formed during sintering,
As a result of preferential solid solution of Cr ₇ C ₃ to iron group metals, only Cr ₂₃ C ₆ is probably present as a lower carbide near the surface of the undissolved Cr ₃ C ₂ phase during or after sintering. It will remain. Another noteworthy aspect of Fukatsu's research results is that he confirms that the lattice constant of the Cr ₃ C _two- phase does not change upon sintering. This fact means that almost no nickel, which is an iron group metal, is dissolved in the Cr ₃ C ₂ phase. Finally, we will discuss the effects of iron group metals that are unavoidably mixed. First of all, it is an extremely well-known fact that the presence of a small amount of iron group metal generally has a significant effect on the densification of a compact made mainly of the refractory hard material. The basis for this is
This is based on the effect as a diffusion promoter in the solid phase reaction stage and the effect as a liquid phase generated by the eutectic reaction. Naturally, the effect of the iron group metal will also be exerted on the green compact of the mixture according to the present invention. However, as has been generally known, this iron group metal is a WC-Co alloy,
As seen in TiC-Ni alloys or Cr ₃ C ₂ -Ni-based alloys, it hardly dissolves into the hard phase, and even after sintering, almost all of it remains as a metal phase at the grain boundaries of the hard phase. For example, it is well understood from the above explanation that the presence of the iron group metal phase has a decisive detrimental effect on the high-temperature properties of the sintered body as a ceramic. However, research on solidified composite materials has revealed that, for example, lower carbides of chromium, ie, Cr ₇ C ₃ and Cr ₂₃ C ₆ , dissolve iron group metals in solid solution. Furthermore, a eutectic reaction exists in the system of a chromium lower carbide containing an iron group metal as a solid solution and an iron group metal-based alloy containing metallic chromium as a solid solution. The eutectic temperature is (Ni・Cr)—(Cr・Ni) ₇ C ₃ , (Co・Cr)—
(Cr・Co) ₇ C ₃ , (Co・Cr) - (Cr・Co) ₂₃ C ₆ systems, 1305℃, 1300℃ and 1340℃, respectively.
It is ℃. These eutectic temperatures are somewhat higher than the eutectic temperature of the Cr ₃ C ₂ -Ni system, which is 1175°C, but
This is a temperature at which it is possible to fully expect the effect of generating a liquid phase and promoting densification in the sintering process. Based on the above considerations, when considering the densification process of a compact consisting of the refractory hard material as the main material, the chromium compound as the binder phase, and a small amount of iron group metal as the auxiliary metal, The following sintering mechanism can be assumed. First, in the solid phase reaction stage at the initial stage of sintering, the orthorhombic chromium compound is solid-dissolved in the same form as the refractory hard material having a more stable and simple crystal form. the result,
A lower-order chromium compound is formed near the surface of the undissolved orthorhombic chromium compound. A part of this lower order compound, for example Cr ₇ C ₃ , is preferentially dissolved in the iron group metal, but at the same time, a part of the iron group metal is also dissolved in the lower order chromium compound. Eventually, the composition of the iron group metal phase containing the lower chromium compound as a solid solution reaches a eutectic composition, and a liquid phase is generated mainly on the surface of the binder phase. As a result, the green compact rapidly becomes densified at a temperature of about 1300°C to 1350°C. At the final stage of densification during the subsequent heating process, or during the holding process at a predetermined sintering temperature, or after densification, a liquid phase consisting mainly of iron group metal-based alloys existing at the binder phase grain boundaries, or a liquid phase in the liquid phase. Lower chromium compounds in which iron group metals remain near the surface of the binder phase, e.g. Cr ₂₃ C ₆
As the iron group metal phase gradually dissolves into the solid solution, the iron group metal phase finally disappears. The dense sintered body obtained based on the above-mentioned sintering mechanism has an iron group metal component mainly dissolved in a low-order compound of chromium, and substantially no metal phase is present. As a result, a dense ceramic based on the refractory hard material with excellent high-temperature properties can be obtained by an inexpensive ordinary sintering method. The second idea of the present invention is the assumption of an orthorhombic chromium compound as a binder phase for the refractory hard material based on the grounds and principles detailed above, and the assumption of a sintering mechanism based on the behavior of this binder phase. It is in. Based on this idea, a mixed powder of one or more of the refractory hard materials and one or more of the chromium compounds was prepared using a ball mill or a vibration mill. Green compacts made of these mixed powders were normally sintered under various conditions, and the effect of the chromium compound as a binder phase was investigated over a wide range of compositions. As a result, green compacts made of these mixed powders generally shrink very easily due to the coexistence of iron group metals that are unavoidably mixed in, and are completely shrunk by ordinary sintering at 1350°C to 1600°C. Or, almost complete densification was confirmed. In addition, most of the chromium compounds in principle form an independent phase as a binder phase in the form of the original orthorhombic compound even after sintering, but as expected, some of them remain as lower-order compounds. was also confirmed. Furthermore, what is most important is that it was confirmed that iron group metals that are inevitably mixed do not exist as a metal phase at all in the dense sintered bodies of any compositions in the examples described below. This proved the validity of the sintering mechanism assumed based on the idea mentioned above. As a result, it has become clear that the dense sintered body obtained by the present invention generally has extremely high hardness and room temperature strength comparable to that of alumina ceramic. In addition, it has become clear that the cutting performance of these dense sintered bodies is equivalent to or better than that of CBN sintered bodies when cutting hardened steel, which is a high-hardness material. Ta. The method for manufacturing the dense sintered body of the present invention will be explained in detail below. First, one or more kinds of the refractory hard material and one or more kinds of the chromium compound having a volume percentage of 1% or more and less than 80% are solidified in the refractory hard material after sintering. A mixed powder is prepared with an amount of iron group metal or iron group metal-based alloy that is below the limit that can exist in a solution state. In this case, the two or more types of refractory hard substances include a combination of two or more of the same compound among carbides, borides, and nitrides, and two or more different types, such as a combination of carbides and borides. Refers to all combinations of two or more types of refractory hard substances including combinations of compounds, and even if the refractory hard substances in these combinations are powders in which all or part of the refractory hard substances have been adjusted to a solid solution in advance. No problem. Also, in the case of the chromium compound, the combination of two or more types refers to all combinations of two or more types, as in the case of the refractory hard material. The powders are mixed by conventional methods, such as wet mixing using a ball mill or a vibratory mill. In this case, the pot and ball materials include cemented carbide,
Cermet, Hastelloy, stainless steel, etc. are used. Mixing time for the former is 20 to 200 hours;
In the latter case, the time is 1 to 10 hours, but in this powder mixing step, iron group metals or iron group metal-based alloys such as cobalt, nickel, and iron are unavoidably mixed in a total amount of about 0.1 to 2 weight percent. Therefore, the iron group metal or iron group metal-based alloy in the above composition does not necessarily need to be blended as a raw material powder, but as a general rule, if the iron group metal or iron group metal-based alloy mixed in the mixing process is used. good. However, if necessary,
For the purpose of adjusting the content of the iron group metal or iron group metal-based alloy, a part thereof may be blended as a raw material powder. As mentioned above, the content of iron group metal or iron group metal-based alloy in the mixed powder is as follows:
In the dense sintered body manufactured from this mixed powder,
The total amount must be adjusted within the limit that allows it to exist as a solid solution in the refractory hard material. If the sintered body contains more iron group metals or iron group metal-based alloys, and some of them exist as a metal phase at the grain boundaries of the refractory hard material phase after sintering, the high-temperature properties of the sintered body are essentially affected. This is because it causes deterioration. In various processes such as powder molding and preliminary sintering, there is no need to take any special considerations to produce the dense sintered body of the present invention, and for example, the production conditions for liquid phase sintered alloys such as cemented carbide and cermets do not require any special considerations. Just follow the instructions. Further, in principle, there is no need to pay special consideration to the sintering conditions, especially the sintering atmosphere. A vacuum atmosphere of 10 ^-3 mmHg to 10 ^-4 mmHg is the most common sintering atmosphere, but considering the vapor pressure of each composition, it is recommended to increase the furnace pressure to about 10 ^-1 mmHg to produce a dense sintered body. It may be effective in obtaining . In addition, for compositions containing nitrides, it is effective to increase the nitrogen partial pressure in the sintering atmosphere in order to prevent the generation of pores due to decomposition and volatilization of nitrides. Although this is a well-known fact, it is of course also effective in producing the dense sintered body of the present invention. Here, in the composition of the dense sintered body of the present invention,
The reason why the proportion of the chromium compound as a binder phase was limited to 1% or more and less than 80% by volume is that if the content of the chromium compound is less than 1% by volume, the sinterability decreases, This is because it is extremely difficult to obtain a dense ceramic sintered body by ordinary sintering, and on the other hand, the properties of a sintered body containing 80% by volume or more approach those of the chromium compound, and the present invention This is because it is difficult to obtain a sintered body with high hardness suitable for the purpose. Note that some details of the densification process and sintering mechanism of each sintered body of the present invention have not yet been elucidated. for example,
In the composition consisting of Mo ₂ C and Cr ₃ C ₂ , the presence of flaky graphite in the sintered body was confirmed, and this fact indicates that Cr ₃ C ₂
This suggests that Cr ₃ C is dissolved in Mo ₂ C as Cr 2 C, but in this case, when Cr ₃ C ₂ is dissolved in Mo ₂ C, low-order Cr 3 C ₂ is formed on the surface of undissolved Cr ₃ C 2. It has not yet been confirmed whether chromium carbide is formed. In addition, although the method for producing a dense sintered body of the present invention is based on a normal sintering method, the method for producing a dense sintered body of the present invention is
Further, it does not impair the significance of the present invention to improve or stabilize the material quality of the sintered body by performing additional pressure sintering using a hot isostatic pressure device. Furthermore, although the dense sintered body of the present invention is primarily intended for use as a cutting tool material, it is possible to easily and inexpensively produce dense sintered bodies with an extremely wide range of compositions, and Since these sintered bodies are high-hardness sintered bodies made of refractory hard materials that inherently have excellent high-temperature properties and corrosion resistance, the dense sintered bodies of the present invention can be used in many fields other than cutting tools. For example, it is expected to exhibit excellent performance in a wide range of applications as heat-resistant materials, corrosion-resistant materials, and wear-resistant materials. Examples of the present invention will be described below. Example 1 Carbide powders of group IVa.Va and VIa metals excluding chromium carbide, namely TiC.ZrC.HfC.VC.
NbC.TaC.Mo2C.WC and _W2C powders each corresponding to 10 weight percent as binder phase _.
Cr ₃ C ₂ powder was blended, acetone was added, and wet ball mill mixing was performed for 72 hours using a pot made of Hastelloy B alloy and a ball made of cemented carbide. For each of these raw material powders, commercially available powders with a particle size of 1 μm to several μm were used. After mixing, acetone was removed by volatilization, and paraffin corresponding to 2% by weight of the powder was added as a molding lubricant. Each mixed powder passed through a 32 mesh sieve contains the iron group metals cobalt, nickel and iron in total.
0.4 to 0.7 weight percent was detected. Each mixed powder was molded into a specimen for a transverse rupture strength test using a mold at a pressure of 1000 Kg/cm ² . These compacts were pre-sintered at 600℃ in a vacuum atmosphere ^of 10 ^-1 mmHg for 1 hour to volatilize and remove the paraffin, and then
1400℃, 1500℃ and Hg vacuum atmosphere
Normal sintering was carried out by heating and holding at each temperature of 1600°C for 2 hours.

【table】

[Table] Table 2 shows the raw material composition of each composition and the relative density and hardness of the obtained sintered bodies. From Table 2, it can be seen that the main material is carbide, which is the refractory hard substance.
Compositions using Cr ₃ C ₂ as a binder phase generally exhibit good sinterability, and can be formed into almost completely dense sintered bodies by ordinary sintering at temperatures of 1600°C or lower. Recognize. In other words, the sinterability differs somewhat depending on each composition, but all compositions have
A relative density of 93.7% or higher was obtained. Furthermore, in the compositions mainly composed of TiC.NbC and Mo ₂ C, which have particularly good sinterability, the relative density already reached 99% or more by ordinary sintering at 1400°C. In X-ray diffraction, each carbonized phase of the main material and a Cr ₃ C ₂ phase were detected as constituent phases of each sintered body.
It was confirmed by microscopic structure observation that no metal phase was present in any of the sintered bodies. Typical microscopic structures are shown in FIGS. 1 to 3. All magnifications are 1000x. Electrolytic corrosion was carried out using a 10% oxalic acid solution. Figure 1 shows an example _of a dense structure in _{a composition} with particularly good sinterability.
This figure shows the structure of a specimen sintered at ℃. Chromium carbide as a binder phase mainly exists at the grain boundaries of the Mo ₂ C phase, but some of it also exists within the grains, contributing to the rapid grain growth of Mo ₂ C particles based on good sinterability. suggest. Although flaky graphite is observed in the structure of this sintered body, it is similar to 90% graphite mainly composed of hexagonal carbide.
The presence of similar flaky graphite was also confirmed in the sintered body of the %W ₂ C - 10% Cr ₃ C ₂ composition. Figure 2 shows 1600, a 90% HfC-10% Cr ₃ C ₂ composition with slightly inferior sinterability.
℃ sintered structure. Chromium carbide is scattered as a rather large binder phase, and at the same time, it is also present in a finely distributed manner at grain boundaries. Figure 3 shows 90% WC - 10% Cr ₃ C ₂
It is a 1600℃ sintered structure of the composition, with extremely fine
The WC phase and chromium carbide phase are uniformly mixed. Furthermore, the grain size of the structure in the sintered body of this composition does not depend on the sintering temperature, and almost no grain growth accompanying high-temperature sintering is observed. Therefore, it is presumed that spinodal decomposition probably occurs in this composition. Example 2 Similar to Example 1, Cr ₃ C ₂ powder corresponding to 10% by weight was blended with TiC powder, and a mixed powder was prepared according to the method of Example 1 using a ball mill. However, in this case, a cermet ball was used, and the mixing time was 96 hours. The green compact of this mixed powder was pre-sintered by heating it at 600°C for 1 hour in a stream of dry hydrogen, and then sintered by heating and holding it at 1400°C for 2 hours in a vacuum atmosphere of 10 ^-1 mmHg. . The obtained sintered body has a relative density of 100%, and the hardness and bending strength are respectively
HRA was 94.7 and 68Kg/ ^mm2 . Moreover, the presence of a metallic phase was not observed in the microscopic structure. Example 3 Eight types of mixed powders were prepared according to the method of Example 1, in which the volume fraction of Cr ₃ C ₂ in the composition of TiC and Cr ₃ C ₂ was varied from 1% to 40%. However, the mixing time was 96 hours. Paraffin was added at 2.5% by weight as a molding lubricant and passed through an 80 mesh sieve. As a comparison sample, a powder prepared by mixing only TiC powder under the same conditions was also prepared. Each mixed powder was molded into a bending test specimen at a pressure of 1000 Kg/cm ² , pre-sintered by heating at 600°C in a stream of dry hydrogen for 1 hour, and then heated at 1400°C.
Sintering was carried out by heating and holding at each temperature of 1500°C and 1600°C for 2 hours. However, in the sintering atmosphere
A vacuum atmosphere of 10 ^-4 mmHg and a hydrogen atmosphere of 5 mmHg were used. In the latter case, each specimen was packed with carbon black and sintered. Table 3 shows each composition.
The Cr ₃ C ₂ content and the relative density, hardness, and bending strength of the obtained sintered body are shown. The second and third specimens showed relative densities slightly over 100%, which is thought to be mainly due to the change in specific volume due to solid solution between components and the accuracy of density measurement. All these relative densities are expressed as 100% in Table 3. By selecting appropriate sintering conditions at a temperature of 1600°C or lower, each of the compositions of the present invention

【table】

【table】

【table】

[Table] A dense sintered body with high hardness of HRA94.0 or higher was obtained. In particular, the amount of Cr ₃ C ₂ as a binder phase is
For compositions in the range of 10 volume percent to 30 volume percent, the properties of the sintered body are generally excellent, with hardness and bending strength of HRA 94.5 or higher and HRA 50, respectively.
A dense sintered body of Kg/mm ² or more was obtained. On the other hand, a dense sintered body could not be obtained with a powder mixed with only TiC powder, and the hardness and bending strength were considerably lower than the sintered body of the present invention. Next, in order to investigate the behavior of iron group metal components and Cr ₃ C ₂ in the composition of the present invention, the distribution of iron group metal components and the composition of chromium carbide in the sintered body were measured using an X-ray microanalyzer. investigated. As a result, it was confirmed that the iron group metal components that are inevitably mixed in during the powder mixing process do not exist as a metal phase in the sintered body, and that most of them are dissolved in solid solution in the lower carbides of chromium. Ta. Figure 4 and Table 4 show an example of the survey results. The specimen was made by packing a composition containing 40% by weight of Cr ₃ C ₂ with carbon black.
It was sintered at 1500°C for 2 hours in a mmHg hydrogen atmosphere. First, in Fig. 4, which examines the distribution of iron group metal components, A is a secondary electron beam image showing the structure of the sintered body, B, C, D, E, and F are titanium in the same field of view as A, respectively. Shows the distribution of chromium, nickel, cobalt and iron. As is clear from this figure, iron group metals are not present at grain boundaries and are mainly dissolved in the chromium-containing phase. However, two types of chromium-containing phases exist: (1) a phase that coexists with a fairly high concentration of iron group metal components, and (2) a phase that does not contain much iron group metal components. Table 4 shows the results of analyzing the components of both phases. From the carbon content of each phase, the phase coexisting with the iron group metal component in (1) is
Cr ₂₃ C ₆ , the phase (2) that does not contain much iron group metal component is
It is determined _to be _Cr3C2 . Example 4 17 types of mixed powders shown in Table 5, in which one or more types of different carbides belonging to the refractory hard substance were added to a composition of TiC and Cr ₃ C ₂ , were prepared in Example 3.
Adjusted accordingly. As shown in Table 5, different types of coal

【table】

[Table] In the composition to which one type of carbide was added, different types of carbides were blended in the form of single carbides, but solid solution carbide powder was used to add two or more types of different types of carbides. In addition, in order to adjust the amount of iron group metal in the mixed powder, some compositions contained a small amount of nickel metal as the raw material powder. Other mixed powder preparation conditions are the same as in Example 3. Each of these mixed powders was molded into a predetermined shape using a mold at a pressure of 1500 kg/ ^cm2 , pre-sintered at 500°C for 1 hour in a stream of dry hydrogen, and then sintered normally under various conditions. . Table 5 shows an example of the sintering conditions for each composition, and the relative density, hardness, and bending strength of each sintered body obtained in that case. In all of the compositions of the present invention, dense sintered bodies with relative densities of 98% or more were obtained. Further, in the microscopic structure, the presence of any metallic phase was not observed at all. Next, in order to investigate the cutting performance of the sintered body for high-hardness work materials, ZrC or HfC was added at 25% by weight as a binder phase.
Two compositions containing 30 weight percent Cr ₃ C ₂ , namely 422C and 425C in Table 5, were subjected to lathe cutting tests. The tools are
The NILR-44 tip shape is SNG432, and a negative brand with an angle of -25° and a width of 0.1 mm is attached to the cutting edge. Commercially available CBN sintered bodies and alumina ceramics were used as comparison samples. All of these comparative samples were tested in commercially available condition. First, the work material is high carbon chromium bearing steel type 3 (SUJ3).
Select a hardened material (hardness: HRC93.5) and set the circumferential speed to 50.
Dry cutting was performed under the following conditions: m/min, depth of cut 0.5 mm, and feed rate 0.1 mm/rotation. In this case, one type of CBN sintered body and one type of ceramic were used as comparative samples. The former uses TiN as a binder, and the latter uses Al ₂ O ₃ -TiC ceramic. The results are shown in Figure 5, and it can be seen that the wear resistance properties of the sintered body of the present invention in cutting high-hardness workpieces are almost equivalent to that of the CBN sintered body for both sintered bodies of both compositions. confirmed. Although the sintered body suffered from slightly larger boundary wear than the CBN sintered body, no chipping was observed at all until the cutting time reached 120 minutes. On the other hand, with Al ₂ O ₃ -TiC ceramic, large chipping occurred after 35 minutes of cutting time, making it impossible to cut. Next, we used a quenched chromium molybdenum steel type 4 (SCM4) material (hardness: HRC54) as the work material, and
A dry cutting test was conducted under the conditions of 100 m/min, depth of cut 0.5 mm, and feed rate 0.1 mm/rotation. The sintered body has
Using 425C in Table 5, four types of comparison samples were used.
A CBN sintered body and two types of alumina ceramics were used. Among the latter ceramics, one type is
Al ₂ O ₃ - TiC ceramic. The wear curve is shown in FIG. It was confirmed that the wear resistance of the sintered body of the present invention is clearly superior to that of conventional alumina ceramics. Also, for CBN sintered bodies,
It was confirmed that it exhibited equivalent or better wear resistance properties. Example 5 Eight types of _mixed powders were prepared by adding one or more different compounds belonging to the refractory hard material to a composition of TiC and _Cr3C2 . In this case, at least one of the different compounds includes a sixth compound.
Nitrides were used as shown in the table. In addition, solute carbide was used as part of the raw material powder. The mixed powder was prepared according to Example 3, but for some compositions, a stainless steel pot was used as the ball mill pot, and the mixing time was 137 hours. Other mixed powder preparation conditions were the same as in Example 3. 3% by weight of paraffin was added as a molding lubricant to each of these mixed powders, and the mixture was molded into a predetermined shape using a mold at a pressure of 1000 kg/cm ² . These green compacts were heated at 600 °C in a hydrogen atmosphere of 10 ^-1 mmHg.
After preliminary sintering by heating to 1 hour at 0.degree. C., normal sintering was performed under various conditions. Table 6 shows an example of the sintering conditions for each composition, and the relative density, hardness, and bending strength of each sintered body obtained in that case. In all compositions, dense sintered bodies with relative densities of 96% or more were obtained. Furthermore, no metal phase was observed in the microscopic structure. Figure 7 shows 543C as an example of the dense structure of the sintered body.
This is a microscopic structure observed at 200x magnification on a lapped specimen. Example 6 Cr ₃ C ₂ is used as the binder phase, nitride or boride belonging to the refractory hard substance is used as the main material, and a different compound belonging to the refractory hard substance is added or not as shown in Table 7 Eight types of mixed powders were prepared. For carbides, nitrides, and borides as dissimilar compounds, respective simple compounds were used. The steps of preparing the mixed powder, molding, and preliminary sintering were all carried out under the same conditions as in Example 5. The pre-sintered bodies were normally sintered under various conditions. Table 7 shows an example of the sintering conditions for each composition, and the relative density, hardness, and bending strength of each sintered body. In all compositions, dense sintered bodies with relative densities of 96% or more were obtained. Further, it was confirmed that no metal phase was present in the microscopic structure of each of these sintered bodies, and that all the iron group metal components mixed in the powder mixing process were dissolved in the refractory hard material. Next, the cutting performance of the sintered body of the present invention was investigated.

[Table] Sintered bodies of two types of compositions, 649C shown in Table 7 and 538C prepared in Example 5, were used as cutting test samples. Both the tool and chip shape are the same as in Example 4. First, high carbon chromium bearing steel 3
The work material is hardened SUJ3 material (hardness: HRC93.5), circumferential speed 50 m/min, depth of cut 0.5 mm, feed 0.1.
Dry cutting was performed under the condition of mm/rotation. The wear curve is shown in FIG. As is clear from a comparison with FIG. 5 of Example 4 where the cutting conditions were the same, both sintered bodies of the present invention exhibit extremely excellent wear resistance properties when cutting high-hardness workpiece materials. In this case, almost no boundary wear occurred. Next, regarding the quenched material (hardness: HRC54) of chromium molybdenum steel type 4 (SCM4),
Dry cutting was performed under the following conditions: min, depth of cut 0.5 mm, and feed rate 0.1 mm/rotation. The wear curve is shown in FIG. It was confirmed that the wear resistance of the sintered body of the present invention was clearly superior to that of the alumina-based ceramic used as a comparative sample, and that it had performance equivalent to or better than that of the CBN sintered body. Furthermore, for the purpose of investigating high-speed cutting performance, 538C
About cast iron FC25 (hardness: HB126~134) as work material, circumferential speed 1000 m/min, depth of cut 1.5 mm, feed.
Dry cutting was performed under the condition of 0.2 mm/rotation. The average flank wear after 5 minutes of cutting was 0.41 mm. On the other hand, the CBN sintered body developed chipping after 1 minute of cutting time, making it impossible to cut. Example 7 TiC with 10 weight percent Cr ₃ C ₂ as binder phase
A mixed powder of a composition containing CrB and 20 weight percent of CrB was prepared under the same conditions as in Example 3. This mixed powder was molded and pre-sintered under the same conditions as in Example 3, and then heated and held at 1600° C. for 4 hours in a vacuum atmosphere of 10 ^-1 mmHg for sintering. The resulting sintered body exhibits a relative density of 100 percent, with hardness and bending strength of HRA 94.2 and 56, respectively.
It was Kg/ ^mm2 . Figure 10 shows a microscopic structure at 1000x magnification after electrolytic corrosion using a 10% oxalic acid solution, and a completely dense structure consisting of two phases, a hard phase and a binder phase, is confirmed.

[Table] Example 8 Five types of mixed powders shown in Table 8 were prepared, which were mainly composed of one or more carbides or borides belonging to the refractory hard substance and mixed with CrB as a binder phase. did. A pot made of Hastelloy B alloy and a ball made of cemented carbide were used to mix the raw material powders, hexane was added, and the mixture was mixed in a wet vibration mill for 3 hours. Paraffin corresponding to 3% by weight was added to each mixed powder as a molding lubricant, and then molded into specimens for transverse rupture strength testing using a mold at a pressure of 1000 kg/cm ² . Each green compact was presintered by heating at 600°C for 1 hour in a dry hydrogen stream with a dew point of -60°C, and then sintered normally under various conditions. Table 8 shows an example of the sintering conditions for each composition and the density, hardness, and bending strength of the obtained sintered bodies. It was confirmed that a dense sintered body could be obtained with any of the compositions. In this case as well, the presence of a metallic phase was not confirmed in the microscopic structure.

[Brief explanation of the drawing]

Figure 1 shows that the composition described in Example 1 is 90%.
This is a microscopic structure magnified 1000 times showing the dense structure of the sintered body of the present invention, which is Mo ₂ C--10% Cr ₃ C ₂ . Figure 2 is similar to Figure 1, and the composition in Example 1 is 90%.
This figure shows a microscopic structure that is 1000 times more dense than that of the sintered body of the present invention, which is HfC-10%Cr ₃ C ₂ . Similarly to Figure 1, Figure 3 also shows that the composition in Example 1 is 90%.
This figure shows a microscopic structure that is 1000 times more dense than the sintered body of the present invention, which is WC-10% Cr ₃ C ₂ . As an example of the distribution of elements in the sintered body of the present invention, FIG. 4 shows the element distribution of the TiC-Cr ₃ C ₂ based sintered body in which the Cr ₃ C ₂ content is 40% by volume in Example 3. In Figure 4, A is a secondary electron beam image showing the structure, and B, C, D, E, and F are the distributions of titanium, chromium, nickel, cobalt, and iron, respectively, in the same field of view as A. . FIG. 5 shows wear curves of the sintered body of the present invention in Example 4 and a comparative sample during dry cutting of hardened high carbon chromium bearing steel (SUJ3) material. 6th
The figure shows the wear curves of the sintered body of the present invention in Example 4 and a comparative sample in dry cutting of hardened chromium molybdenum steel (SCM4) material. FIG. 7 shows an example of the dense structure of the sintered body of the present invention in Example 5, where the composition is 35%TiC-25%HfC-
This figure shows the microscopic structure of the sintered body of the present invention, which is 8% NbC-2% TiN-30% Cr ₃ C ₂ , in a no-etch state at a magnification of 200 times. FIG. 8 shows wear curves in dry cutting of high carbon chromium bearing steel (SUJ3) hardened materials of the sintered bodies of the present invention in Examples 5 and 6. FIG. 9 shows the wear curves of the sintered body of the present invention and the comparative sample in dry cutting of hardened chromium molybdenum steel (SCM4) material in Examples 5 and 6. Figure 10 shows _that the composition of Example 7 is 70%TiC-10% _Cr3C2-
In the sintered body of the present invention containing 20% CrB, the dense structure consists of two phases: a hard phase and a binder phase.
This is a microscopic structure magnified 1000 times.

Claims (1)

[Claims] 1 Carbides of group IVa and Va group metals of the periodic table,
The main component is powder of one or more single compounds of group VIa metal carbides and borides, excluding borides, nitrides, and chromium compounds, or solid solution powders thereof as the hard phase, and as the binder phase. a single compound powder or a solid solution powder of one or more of orthorhombic chromium carbide and orthorhombic chromium boride having a volume percentage of 1% or more and 80% or less, and the bond after sintering. A compact of a powder mixture consisting of one or more iron group metal powders or iron group metal-based alloy powders as auxiliary metals in an amount within the limit that can be present in solid solution in the phase and the hard phase. , dense and high hardness characterized by the absence of a binder metal phase, obtained by normal sintering under conditions where the entire amount of auxiliary metal exists in a solid solution state in the binder phase and hard phase after sintering. Ceramic sintered body.