JPH0372696B2 - - Google Patents

Description

[Detailed description of the invention]

(Industrial Application Field) The present invention provides a coated surface-treated sintered alloy which is formed by forming a coating film on the surface of a surface-treated sintered alloy that is effective as a coated sintered alloy used as a cutting tool member or a wear-resistant tool member. This relates to quality sintered alloys. (Prior technology) So-called coated sintered alloys, which are made by coating sintered alloys such as cemented carbide with a thin layer of more wear-resistant TiC, TiCN, TiN, Al ₂ O ₃ , etc., are It combines the toughness of the alloy and the excellent wear resistance of the coating film, and is already in widespread practical use. Although the above-mentioned coating film has excellent wear resistance, it is extremely hard and brittle, so it is easy to crack during use, which causes the crack to spread to the base material and lead to chipping of the cutting edge. There was a problem. Japanese Patent Application Laid-Open No. 54-87719 is an excellent prior art proposed to solve this problem.
It is already in practical use. This prior art is a cemented carbide that has a nitrogen-containing B-1 type crystal structure phase (hereinafter referred to as β phase) and a WC phase as hard phases, and an iron group metal as a binder phase. 5 to 5 of the cemented carbide
The amount of β phase in the 200 μm surface layer is reduced compared to other parts. Cemented carbide with such a special structure is produced by vacuum sintering a green compact consisting of B-1 carbonitride, which essentially contains nitrogen, WC, and iron group metals, and then denitrifying the surface. It is said that it is manufactured by In fact, the phenomenon of deβ removal of the surface layer by this method is
It was studied in detail by Professor Hisashi Suzuki of the University of Tokyo at the time (Journal of the Japan Institute of Metals, Vol. 45, No. 1, pp. 95-99).
Page and "Powders and Powder Metallurgy" Volume 29 No. 2 No. 20
(pages 23 to 23) De-β removal of the surface layer has been confirmed,
As mentioned above, the beta-free cemented carbide is used as a base material for coated cemented carbide. (Problems to be Solved by the Invention) However, it is known that when the surface layer of the cemented carbide obtained by this method is used as a base material for a coated cemented carbide, the fracture resistance is still insufficient. Ta. In other words, Figure 1 is taken from the figure on page 302 of "Cemented Carbide and Sintered Hard Materials" (Maruzen) edited by Hisashi Suzuki. However, when the amount of Co, which is the binder phase, is large, it is possible to suppress the propagation of cracks from the surface and improve fracture resistance. , Co is enriched within several tens of micrometers from the surface, and the surface is rather depleted of Co, with other examples showing that the average concentration inside is at the same level or even lower. is known. Therefore, as a natural consequence, if such a surface layer de-β cemented carbide is used as the base material of the coated cemented carbide, the propagation of cracks that occur in the brittle coating film into the base material can be prevented by preventing the cracks from propagating into the base material surface layer. The deterrent effect will be diminished. The purpose of the present invention is to solve the technical problems of the above-mentioned prior art, and to provide a coated surface-tempered sintered alloy formed by forming a hard film on the surface of a base material having a novel structure useful for coated sintered alloys. It is about providing. (Means for Solving the Problems) The present invention provides at least one of carbides, carbonitrides, carbonates, carbonitrides, or mutual solid solutions of metals from groups 4a, 5a, and 6a of the periodic table. In the coated sintered alloy, a coating film is formed on the surface of the sintered alloy, which is composed of 80 to 97wt% of a hard phase containing iron group metals, and the remainder is a binder phase containing at least one type of iron group metal. The sintered alloy is made of a surface-treated sintered alloy in which the relative concentration of the binder phase of the sintered alloy reaches its maximum value at the surface of the sintered alloy and then decreases toward the inside from the surface of the sintered alloy to at least 5 μm. The coating film contains carbides, nitrides, carbonates, nitrides and mutual solid solutions of metals from groups 4a, 5a and 6a of the periodic table, silicon carbide, silicon nitride, aluminum oxide, aluminum oxynitride, aluminum nitride, cubic This is a surface-treated sintered alloy characterized by being made of a single layer or multiple layers of at least one of crystalline boron nitride and diamond. The surface-treated sintered alloy of the present invention will be explained in more detail using FIG. The relative concentration of the binder phase decreases from a to b toward the internal point b up to at least 5 μm, and after the relative concentration of the bonded phase further inside changes as shown in Fig. 2 or as shown in Fig. 2, the average internal Examples include a case in which the concentration of the bonded phase reaches point c, or in an extreme example, a case in which the relative concentration of the bonded phase decreases from a to b and then immediately reaches point c. Among these, in the layer from the surface of the sintered alloy to the inner average concentration of the binder phase, the so-called surface layer where the relative concentration of the binder phase of the sintered alloy fluctuates, as shown in Fig. 3, for example. After the relative concentration of the binder phase reaches its maximum value (a) at the surface of the sintered alloy, it smoothly decreases from a point (b) at least 5 μm inside the sintered alloy to the average binder phase concentration inside (c). , or after the relative concentration of the binder phase has reached its maximum value (a) at the surface of the sintered alloy, as shown in Figure 4, for example, at an internal point at least 5 μm from the surface of the sintered alloy. It further decreases from (b) to a minimum value (d) that is smaller than the average bonded phase concentration inside, and then increases smoothly toward the inside until it reaches the average bonded phase concentration inside (c). The case is of practical importance. FIG. 5 shows, as an example, the relative concentration distribution of each element from the surface to the inside of the sintered alloy, assuming that the average concentration in the interior of the constituent elements of the surface-treated sintered alloy of the present invention is 1. It is something. That is, as is clear from FIG. 5, the concentration of Co, which is the binder phase, on the surface of the surface-treated sintered alloy of the present invention reaches a maximum value that is higher than the average Co concentration inside, and then decreases. However, it reaches a minimum value that is smaller than the average Co concentration inside, and then increases until it reaches the average Co concentration inside. The hard phase of the surface-treated sintered alloy in the present invention is
If the amount exceeds 97% by weight, the binder phase becomes relatively less than 3% by weight, resulting in a significant decrease in strength as a coated surface-tempered sintered alloy, and conversely, the hard phase of the surface-tempered sintered alloy If it is less than 80% by weight, the binder phase will be relatively large, exceeding 20% by weight, and the life of the coated surface-tempered sintered alloy will be significantly reduced due to its plastic deformation resistance and wear resistance. , 80～
It was set at 97% by weight. In addition, if the relative concentration of the binder phase in the surface-treated sintered alloy of the present invention reaches its maximum value at the surface and then decreases toward the inside within a distance of less than 5 μm from this surface, the maximum binder phase amount at the surface Although it is related to
It was determined that the thickness decreases inward to 5 μm. Next, as a method for manufacturing the above-mentioned surface-treated sintered alloy in the present invention, 4a, 5a and 5a of the periodic table are used.
Consisting of a hard phase containing at least one of group 6a metal carbides, carbonitrides, carbonates, carbonitrides, or mutual solid solutions thereof, and a binder phase containing at least one iron group metal. During the manufacturing process of the sintered alloy, after or during the sintering process, the surface of the sintered alloy is decarburized at a temperature within the solid-liquid coexistence temperature range of the binder phase. A sintered alloy can be obtained in which the relative concentration reaches a maximum value at the surface of the sintered alloy, and then decreases toward the inside by at least 5 μm from the surface of the sintered alloy.
At this time, preferably after sintering the sintered alloy or during the sintering process, the surface of the sintered alloy is decarburized at a slow rate at a temperature within the solid-liquid coexistence temperature range of the binder phase. A surface-treated sintered bond in which the relative concentration of the bonding phase of the bonded gold is maximum at the surface of the sintered alloy, and decreases toward the inside for at least 5 μm from the surface, reaching the average concentration of the bonding phase inside. obtaining gold, and preferably decarburizing the surface of the sintered alloy at a rapid rate or decarburizing the surface of the sintered alloy after carburizing the surface of the sintered alloy. As a result, the relative concentration of the binder phase of the sintered alloy is maximum at the surface of the sintered alloy, and at least 5 μm from the surface.
The surface-treated sintered bond decreases toward the inside until it reaches a minimum value smaller than the average binder phase concentration inside, and then increases smoothly toward the inside until it reaches the average binder phase concentration inside. It's something you can earn money from. Furthermore, the present invention uses the above-mentioned surface-treated sintered alloy as a base material, and the surface of the base material has, for example,
Carbides, nitrides, carbonates, nitrides and their mutual solid solutions of metals of groups 4a, 5a and 6a, silicon carbide, silicon nitride, aluminum oxide, aluminum oxynitride, aluminum nitride, cubic boron nitride, in diamond. The present invention provides a coated surface-tempered sintered alloy formed with at least one type of single-layer or multilayer coating film. This coated surface-tempered sintered alloy can be formed by forming a coating film on the surface of the above-mentioned surface-tempered sintered alloy using the conventional chemical vapor deposition method (CVD method) or physical vapor deposition method (PVD method). You can get it by doing this. Of course, a surface layer formed by combining a conventional surface layer with a β-free phase and a surface layer with the above-mentioned binder phase concentration distribution,
Specifically, a state in which a hard phase with a B-1 type crystal structure does not exist in the surface layer, or even if it does exist, it is in a state in which it is smaller than in the interior of the sintered alloy excluding the surface layer, or in other words, The surface layer is a phase in which the hard phase of the B-1 type crystal structure is reduced compared to the inside of the sintered alloy excluding the surface layer, that is, a non-β layer, and the surface layer has the above-mentioned binder phase concentration distribution. It is also preferable to form a coating film on the surface of the sintered alloy having the following properties to further improve the fracture resistance of the surface layer. In this case, not only the distribution of the binder phase shown in Fig. 3 but also the distribution of the binder phase shown in Fig. 4 can be easily obtained, and the coated sintered alloy can be used as a cutting tool member. When used, it has the advantage of improving the strength of the cutting edge. The metals of groups 4a, 5a, and 6a of the periodic table mentioned here are Ti, Zr, Hf (group 4a metals), V, Nb, and Ta.
(Group 5a metals) and Cr, Mo, W (Group 6a metals). It represents the average concentration of the bonded phase within a stable region within the bonded gold where the relative concentration of the bonded phase hardly changes. Expressed another way,
It represents the average concentration of the binder phase inside the sintered alloy excluding the surface layer. The depth of the surface layer of the sintered alloy varies depending on the shape of the sintered alloy and the conditions for surface treatment. Both can be controlled over a wide range, such as when the diameter is 2 mm. Practically speaking, it is preferable to determine the shape of the sintered alloy, the material and structure of the coating film, and the intended use. (Function) The present invention provides a sintered alloy that essentially contains carbon, and the surface of the sintered alloy is decarburized by reheating the sintered alloy to the solid-liquid coexistence temperature of the binder phase in a decarburizing atmosphere. This was based on the knowledge that only the binder phase can be enriched in the surface layer by using the method. Although the mechanism by which this bonded phase enrichment phenomenon occurs is not necessarily clear, it is thought to be based on the following principle. For convenience of explanation, the simple W-
This will be explained using a cross-sectional phase diagram at 16% Co in the Co-C ternary phase diagram. The sintered alloy that is the raw material may be produced by any known method, but the sintered alloy produced in this way has a solid-liquid coexistence temperature of the binder phase shown in the shaded area in Figure 6. heated to a temperature within the range. At this time, if the atmosphere in the furnace is made into a decarburizing atmosphere by, for example, introducing CO ₂ gas or creating a high vacuum, decarburization will occur on the surface of the sintered alloy, and the carbon concentration on the surface will change to the shaded area in Figure 6. From the liquidus line AB of the bonded phase to the solidus line of the bonded phase in the enlarged figure 7
It decreases in the direction of CD, as shown by arrow b, and reaches the solidus line CD, where the liquid phase solidifies and a volume decrease occurs accordingly. As a result, the liquid phase is replenished from inside, but this also reaches near the surface and is decarburized, reaching the solidus CD
The process is repeated until the surface is enriched with the binder phase. The reason why the concentration of the bonded phase becomes smaller as shown in Figure 1 on the surface after β removal as disclosed in the prior art Japanese Patent Application Laid-Open No. 54-87719 is explained in the research by Professor Suzuki et al. (Vol. 45, p. 98)
According to the above, it is thought that this is due to evaporative diffusion during sintering, but in the case of the present invention, there is no evaporative diffusion, and as a result, the maximum concentration at the surface can be maintained because of the solidification of the surface bonding phase due to surface decarburization. This is thought to be due to the fact that evaporation could be avoided. Now, when producing the surface-treated sintered alloy of the present invention, the liquid phase supplied to the surface is naturally procured fastest from a portion relatively close to the surface, so if the decarburization treatment is performed quickly, the liquid phase is A shortage of liquid phase occurs in the vicinity, creating a minimum point of bonded phase concentration.
On the other hand, if the decarburization process is performed slowly, this minimum point can be virtually eliminated. For example, H ₂ +
When decarburizing WC-5%Co alloy using CO ₂ atmospheric gas, the CO ₂ gas concentration in the atmospheric gas is 10% or more, the atmospheric gas pressure is 10 torr or more, the temperature is 1330°C or less,
If the treatment time is within 3 minutes, it is a rapid decarburization treatment and can create a minimum value in the relative concentration distribution of the binder phase. On the other hand, the CO ₂ concentration in the atmospheric gas is 10% or less,
If the atmospheric gas pressure is 10 torr or less, the temperature is 1330° C. or more, and the treatment time is 3 minutes or more, decarburization is slow and there is substantially no minimum value in the relative concentration distribution of the binder phase.
Furthermore, after sintering or during the sintering process, the solid-liquid coexistence temperature range of the binder phase can be passed through while slowly cooling at a rate of 10° C./min or less. At this time, the cooling rate does not necessarily have to be constant; for example, 1270℃
It is also possible to change the speed during slow cooling, such as setting the speed to 1°C/min until then and then increasing the speed to 3°C/min.
It is advantageous to control the depth and microscopic texture of the surface layer. Further, in the decarburization treatment, the above-mentioned decarburizing gas may be used as the atmospheric gas, or the atmosphere may be a vacuum if the carbon content of the sintered alloy is high. In general, sintered alloys with high binder phase content,
Alternatively, in a sintered alloy with a high C content, the enrichment phenomenon on the surface of the binder phase occurs quickly due to the decarburization treatment, so each of the above conditions is appropriately controlled depending on the sintered alloy used. That is, in addition to the strength, temperature, and time of the decarburizing atmosphere, it is also necessary to control the cooling rate, since in slow cooling, the smaller the cooling rate, the greater the amount of binder phase enrichment in the surface layer. Note that if this decarburization operation is carried out particularly rapidly in a stronger decarburization atmosphere, the binder phase and the hard phase appear in alternating layers parallel to the surface in the binder phase enriched region near the surface. Furthermore, the surface-treated sintered alloy of the present invention can be used in various alloy systems containing carbon, including the simplest cemented carbide of the WC-Co system, as well as carbide-based cermets that do not contain nitrogen, and of course contain nitrogen. This is an epoch-making technology that can also be applied to cermets. As can be seen from the above, when producing the surface-treated sintered alloy of the present invention, the decarburization operation does not necessarily need to be performed after sintering. In other words, during the sintering process, the temperature may once fall below the solid-liquid coexistence temperature of the binder phase and then be raised again to the solid-liquid coexistence temperature of the binder phase to decarburize, or during the sintering process, the temperature The surface-treated sintered alloy of the present invention can be obtained even if decarburized at a solid-liquid coexistence temperature of . On the other hand, if the surface of the sintered alloy is carburized at the solid-liquid coexistence temperature of the binder phase, the amount of carbon on the surface increases in the direction opposite to the arrow b in Fig. 7, reaching the binder phase liquidus line AB, resulting in the above-mentioned phenomenon. The opposite phenomenon occurs. If the above-described decarburization operation is performed after such carburizing treatment, the valley at the minimum binder phase concentration can be made even deeper and stably produced. The carburizing treatment can be carried out, for example, in a reduced pressure mixed gas of H ₂ +CH ₄ . At this time, since the amount of binder phase enrichment in the surface layer varies depending on the carbon content in the alloy before decarburization, the carburizing treatment conditions are also appropriately controlled. Furthermore, by performing the carburizing treatment before the decarburization treatment, the binder phase concentration distribution once reached a minimum value as described above, then increased, and again exceeded the average internal binder phase concentration and reached a small maximum value. Although the internal bonded phase concentration may later become average, it does not pose any problem. Note that setting a minimum value in the binder phase concentration distribution of the surface layer has the effect of suppressing plastic deformation, especially in applications where a large amount of heat is generated. Further, in the present invention, in which the surface of the surface-treated sintered alloy obtained in this manner is coated with a coating film having excellent wear resistance, the surface layer of the surface-treated sintered alloy is It compensates for the weak point of brittleness and allows the coating film to fully exhibit its excellent wear resistance. (Example) Example 1 Commercially available WC, WC-TiC solid solution (WC/TiC=70/30 weight ratio), TaC and
Using various Co powders (particle size 1.5 to 3 μm), 85% WC
After blending to have a composition of -5% TiC - 3% TaC - 7% Co (wt%), wet ball milling was performed for 48 hours using acetone as a solvent. After the mill,
After drying, it was press-molded into a JIS transverse rupture strength piece shape, and then sintered in a vacuum at 1400°C for 1 hour. After surface polishing, these were divided into two parts, and one group was heated at 1330℃ for 10 minutes in 80% H ₂ -20% CH ₄
After carburizing in a mixed gas of 20 torr, decarburization was performed at 1310° C. for 2 minutes in a mixed gas of 90% H ₂ -10% CO ₂ of 10 torr, and then furnace cooling was performed in a vacuum. For these samples, W, Ti,
EPMA analysis of the elemental concentration distribution of Ta and Co as a function of depth from the surface yielded the results shown in Figure 5. Here, the concentration of each element is normalized and shown with each concentration at the center of the sample being 1. Than this,
The amount of Co reaches its maximum at the surface of the sample, decreases continuously toward the inside, reaches its minimum value, and then reaches its internal value. The amounts of W, Ti, and Ta showed opposite trends in response to changes in the amount of Co. On the other hand, for the untreated sample, the concentrations of each element both inside and outside the sample showed constant values. Next, the treated sample and the untreated sample were coated with TiC to a thickness of 5 μm by chemical vapor deposition. Then, when we measured the transverse rupture strength according to JIS,
The untreated sample had an average strength of 145 Kg/mm ² , while the surface-treated sample had a high strength of 205 Kg/mm ² , an average of 20 samples. Example 2 Using various commercially available raw material powders, a blending composition of 86% WC - 4% TiC - 3% TaC was prepared according to a conventional manufacturing method.
A plurality of SNMN 432-shaped compacts containing -1% NbC - 6% Co (wt%) were prepared. Some of these were heated at 1200℃ for 30 minutes during the temperature increase during sintering.
After nitriding in nitrogen gas at 30 torr, sintering was performed in vacuum at 1420°C for 1 hour. All remaining compacts were vacuum sintered at 1420°C for 1 hour without undergoing a nitriding process. The treatments shown in Table 1 were carried out with some exceptions. After treatment, the distribution of Co content in the vertical cross section of each sample was determined as a function of depth from the surface.
EPMA analysis showed that the central value of the sample was 100%
The distribution shown in Table 1 was confirmed. Next, all samples were coated with TiC1μ by chemical vapor deposition.
A coated cemented carbide was obtained by sequentially coating with 1 μm of TiCN, 4 μm of TiCN, and 1 μm of Al ₂ O ₃ . These were subjected to a fracture resistance test and a wear resistance test by peripheral turning under the conditions shown below, and the results shown in FIGS. 8 and 9, respectively, were obtained. (1) Fracture resistance test workpiece S48C (H _B 255) with equally spaced slots. Cutting speed 100m/min Depth of cut 1.5mm Feed rate 0.3mm/rotation No cutting oil (dry cutting) (2) Wear resistance test Work material S48C (H _B 240) Cutting speed 180m/min Depth of cut 1.5mm Feed rate 0.24 mm/rotation No cutting oil (dry cutting)

【table】

[Table] From the above results, it can be seen that the samples of the present invention have excellent properties in which the fracture resistance is greatly improved without decreasing the wear resistance, and the variation in fracture resistance is extremely small. Example 3 Following the same manufacturing method as Example 2, the composition was 88%
A plurality of TNMN332-shaped samples of WC-2%TiC-4%TaC-5%Co-1%Ni (wt%) were all produced by vacuum sintering at 1400°C for 1 hour. These samples were then divided into three parts and subjected to surface treatment under the conditions shown in Table 2. In the vertical section of the sample
Distribution of Co+Ni content as a function of depth from the surface
The results of EPMA analysis are calculated as 100% of the central value of the sample.
as shown in the table. Next, these samples were coated with TiC2μ by chemical vapor deposition.
2 μm of TiCN and 2 μm of TiN were sequentially coated. Then, a fracture resistance test was conducted under the same conditions as in Example 2, and the results shown in FIG. 10 were obtained. This shows that the distribution of the binder phase on the surface has a large effect on the variation in fracture resistance, and that the fracture resistance is extremely stable when the amount of binder phase is maximum on the surface.

[Table] Example 4 According to the same manufacturing method as Examples 2 and 3, the composition was
86%WC-3%TiC-1%TiN-3%TaC-7%
A plurality of TNMG332-shaped powder compacts of Co (wt%) were prepared. Among these, the samples of the present invention shown in Table 3 and the green compacts used in the samples were heated at 1420°C in vacuum.
After sintering by holding for 1 hour, the treatments shown in Table 3 were performed during the cooling process in the same furnace. The compacts used in the comparative samples shown in Table 3 were sintered under the same conditions as described above and immediately cooled. The amount of Co in the vertical section of each sample thus obtained was analyzed by EPMA as a function of depth from the surface, and the results are shown in Table 3, with the Co value at the center (inside) of the sample as 100%. Ta. In addition, for all samples, from the surface
It consists of WC-Co phase up to a depth of 20 μm, and B-1
The hard phase of the type crystal structure had disappeared. Then, all samples were coated with TiC4μ by chemical vapor deposition.
A coated cemented carbide was obtained by sequentially coating with 2 μm of Al ₂ O ₃ and 2 μm of Al 2 O 3 . The coated cemented carbide obtained in this way was subjected to a fracture resistance test under the conditions shown in Example 2, in which only the feed rate was changed, i.e., the feed rate started at 0.15 mm/rotation, and the number of impacts on the cutting edge was 4000 at the same feed rate. If no breakage occurred even after reaching the number of rotations, the feed amount was increased by 0.05 mm/rotation, and a test was conducted to determine superiority or inferiority based on the feed amount when the breakage occurred.
The feed rate at the time of breakage was low at 0.2 to 0.3 mm/rotation, and there was some variation, whereas the sample of the present invention showed small variation and broke at a feed rate of 0.35 to 0.4 mm/rotation, and the sample of the present invention No breakage occurred even at the maximum feed rate of 0.4 mm/rotation. On the other hand, using S48C (H _B 230) round bar as the work material, dry continuous turning was performed for 10 minutes at a cutting speed of 200 m/min, depth of cut of 1.5 mm, and feed rate of 0.3 mm/rotation.
When the amount of plastic deformation of the cutting edge of each sample was determined, it was 0.03 mm for the comparative sample, 0.03 mm for the sample of the present invention, and 0.04 mm for the sample of the present invention. As a result, it was found that the coated cemented carbide of the present invention has substantially the same plastic deformation resistance and significantly superior fracture resistance as compared to the comparative coated cemented carbide.

[Table] Example 5 The various commercially available raw material powders used in Examples 1 and 2 were mixed to have the composition shown in Table 4, and then sintered in the same manner as in Example 1. got money. (However, only the sintering temperature was 1420°C for inventive product 12 and comparative products 17 and 18, 1400°C for inventive product 13, 14, 15 and comparative products 19 and 20, and 1400°C for inventive product 16 and comparative products 21,
22 is 1360℃. ) Among the sintered alloys obtained in this way, inventive products 12 to 16 and comparative products 17, 20, and 22 were carburized and decarburized under the conditions shown in Table 4, and other comparative products
18, 19, and 21 were left untreated. Next, a coating film similar to that in Example 2 was formed on the surface of each of the sintered alloys of Inventive Products 12 to 16 and Comparative Products 17 to 22, and then the fracture resistance test and resistance test conducted in Example 2 were performed. A wear test (cutting time 20 minutes) was conducted and the results are shown in Table 5. In addition, the surface parts of the sintered alloys of the present invention products 12 to 16 and comparative products 17 to 22 were examined in the same manner as in Example 2,
The surface of each sintered alloy when the binder phase concentration inside each sintered alloy is 100.
The depth from the surface at which the amount of Co and the concentration of the binder phase were minimum were determined and listed in Table 5. From the results in Table 5, it can be seen that products 12 to 16 of the present invention are superior in both wear resistance and chipping resistance in a well-balanced manner compared to comparative products deviating from the present invention and conventional comparative products 17 to 22. It is.

【table】

[Table] (Effects) As mentioned above, in the surface-treated sintered alloy of the present invention, the relative concentration of the binder phase is substantially at its maximum at the surface, so that cracks that occur in the brittle coating film can be removed from the base material. Progress is blocked on the surface, which has the effect of preventing tool breakage. Furthermore, if necessary, it is possible to create a portion at an appropriate depth from the surface where the binder phase concentration of the surface-treated sintered alloy of the present invention is smaller than the average binder phase concentration inside. The part with the highest concentration of binder phase on the surface not only prevents cracks from propagating into the interior, but also the part with the lowest binder phase concentration firmly prevents plastic deformation of the cutting edge, which is often a problem in high-speed heavy cutting. This has the effect of preventing plastic deformation and resulting damage to the cutting edge.

[Brief explanation of drawings]

FIG. 1 shows the relative concentration distribution of Co, W, and Ti from the surface to the inside of a conventional sintered alloy. FIG. 2, FIG. 3, and FIG. 4 show typical examples of the relative concentration distribution of the binder phase of the surface-treated sintered alloy according to the present invention. FIG. 5 shows the relative concentration distribution of Co, W, Ti, and Ta in the surface-treated sintered alloy of the present invention determined in Example 1. FIG. 6 shows a cross-sectional phase diagram of WC-16wt%Co. FIG. 7 is an enlarged view of the solid-liquid coexistence region of the bonded phase in FIG. 6, where AB indicates the liquidus line of the bonded phase and CD indicates the solidus line of the bonded phase. Figure 8 is a graph of the fracture resistance test results in Example 2;
The figure shows a graph of the wear resistance test results in Example 2, and FIG. 10 shows the graph of the fracture resistance test results in Example 3.

Claims (1)

[Scope of Claims] 1. Carbide of group 4a, 5a and 6a metals of the periodic table,
Hard phase containing at least one of carbonitrides, carbonates, carbonitrides, or mutual solid solutions thereof80~
97wt% and the remainder at least one of the iron group metals
In a coated sintered alloy formed by forming a coating film on the surface of a sintered alloy consisting of a binder phase containing seeds, the relative concentration of the binder phase of the sintered alloy reaches its maximum value on the surface of the sintered alloy. After that, at least
The coating film consists of a surface-treated sintered alloy that decreases inward to 5 μm, and the coating film is 4a of the periodic table.
carbides, nitrides, carbonates of group 5a and 6a metals;
Nitrogen oxides and their mutual solid solutions, silicon carbide,
A surface-treated sintered alloy comprising a single layer or multiple layers of at least one of silicon nitride, aluminum oxide, aluminum oxynitride, aluminum nitride, cubic boron nitride, and diamond. 2. Claim 1, characterized in that the relative concentration of the bonding layer smoothly approaches the average bonding phase concentration inside the sintered alloy within at least 5 μm from the surface of the sintered alloy. coated surface tempered sintered alloy. 3 After the relative concentration of the binder phase reaches a minimum value smaller than the average binder phase concentration inside at least 5 μm from the surface of the sintered alloy, the average binder phase concentration inside the sintered alloy increases further inward. The coated surface-tempered sintered alloy according to claim 1, characterized in that the coated surface-tempered sintered alloy has a smooth increase in temperature. 4 carbides of metals from groups 4a, 5a and 6a of the periodic table;
Hard phase containing at least one of carbonitrides, carbonates, carbonitrides, or mutual solid solutions thereof80~
97wt% and the remainder at least one of the iron group metals
In a coated sintered alloy formed by forming a coating film on the surface of a sintered alloy consisting of a binder phase containing seeds, the relative concentration of the binder phase of the sintered alloy reaches its maximum value on the surface of the sintered alloy. After that, at least
The thickness decreases inward to 5 μm, and the hard phase of the sintered alloy is tungsten carbide and 4a and 5a of the periodic table.
and carbides, carbonitrides, carbonates of group 6a metals,
and at least one B-1 type crystal structure phase among carbonitride oxides or mutual solid solutions thereof, and the B-1 type in the inner surface layer up to at least 5 μm from the surface of the sintered alloy. It is made of a surface-treated sintered alloy in which the hard phase of the crystal structure is reduced compared to the hard phase of the B-1 type crystal structure in the interior excluding the surface layer, and the coating film is 4a of the periodic table. ,5a
and at least one of group 6a metal carbides, nitrides, carbonates, nitrides, and mutual solid solutions thereof, silicon carbide, silicon nitride, aluminum oxide, aluminum oxynitride, aluminum nitride, cubic boron nitride, and diamond. A coated surface-tempered sintered alloy characterized by comprising a single layer or multiple layers of.