JP4911937B2 - High-strength cemented carbide, manufacturing method thereof and tool using the same - Google Patents

Description

  The present invention relates to a WC-Co-based high strength and high toughness cemented carbide, and is excellent in wear resistance, toughness, fracture resistance, and heat cracking resistance, cold forging tools, rolls, bits for mining tools Applicable to crushing blades, cutting blades, and other wear-resistant tools. The WC-Co system in the present invention includes not only hard particles mainly composed of WC and iron group metal powders containing Co, but also excludes WC of the periodic table IVa, Va, VIa group elements as hard particles. Meaning at least one selected from carbides, nitrides, carbonitrides and borides.

  A commercially available cemented carbide for wear resistance is a composite material of a WC hard phase and a Co metal phase, and is a typical dispersion alloy. Its mechanical properties depend on the particle size of the WC hard phase and the amount of Co-bonded metal phase, and in particular, hardness and toughness are in a trade-off relationship. In order to make full use of its extremely excellent hardness, many proposals have been made regarding cemented carbide with high strength and high toughness.

  For example, Japanese Examined Patent Publication No. 47-23049 discloses a high-strength alloy composed of tungsten carbide plate-like particles having unequal dimensions and a maximum dimension of 50 μm or less and at least three times the minimum dimension and an Fe group metal. It is shown. However, plate-like tungsten carbide of unequal dimensions is intended to obtain an oriented WC grain growth structure by applying a shearing force by rolling while using fine tungsten carbide as a starting material while heating. Various wear-resistant carbide products that require near-net product shapes have a problem that they are difficult to apply.

  Furthermore, Japanese Patent Laid-Open No. 02-274827 relates to a technique for producing an anisotropic cemented carbide molded body having excellent crack propagation resistance or toughness, and the sintered cemented carbide is oxidized and reduced. A method for obtaining a mixed powder of WC and Co that is carbonized and has anisotropy is described later, but this is a method of reusing used cemented carbide and requires special equipment. difficult.

  These inventions adopt a unique particle form such as anisotropic WC particles or plate-like tungsten carbide as the hard phase, thereby making it possible to produce a high-hardness and high-toughness cemented carbide with a uniform structure throughout the product. It is a manufacturing method. On the other hand, a method for producing a high-strength cemented carbide as a composite material has also been proposed.

  That is, in Japanese Patent Application Laid-Open No. 08-127807, a graded composite material in which a grain growth promoting material is impregnated from the surface of a molded body and main firing is performed after drying, whereby the surface layer portion becomes a ceramic grain growth structure and the inside becomes a metal phase rich. It is shown.

  Furthermore, JP-A-2002-249843 discloses that a mixed powder of oxide ceramic particles and metal particles is molded to form a molded body, and a boron compound-containing solution is applied to the surface of the molded body and sintered to obtain a surface layer. It is shown that a composite material having both a high hardness, a high strength and a high toughness having a grain growth structure and a three-dimensional network structure in the part can be obtained. However, these proposals are only toughening by the grain growth structure of the surface layer portion, and do not mention, for example, that the particle size of the surface layer portion is smaller than that of the inner quality portion.

On the other hand, in Japanese Patent Laid-Open No. 04-128330, in a sintered alloy having a hard layer mainly composed of metal carbide and a binding layer composed of iron-based metal, various diffusion elements are applied to the surface of the pressure-formed body before sintering. Then, the diffusion element and the bonding layer are reacted on the surface of the hard phase by liquid phase sintering, and the concentration of the bonding phase gradually increases from the surface to the inside, and the average particle size of the hard phase gradually increases. Sintered alloys having a graded composition have been proposed.
Japanese Patent Publication No. 47-23049 Japanese Patent Laid-Open No. 02-274827 Japanese Unexamined Patent Publication No. 08-127807 JP 2002-249843 Japanese Patent Laid-Open No. 04-128330

Cutting and turning tips, which are the main uses of cemented carbide, are determined by die molding, so the application of the above-mentioned plate-like WC and anisotropic WC is extremely easy, but they have complicated shapes. However, application to wear-resistant cemented carbide products manufactured by various forming processes is extremely difficult. In addition, the conventionally proposed sintered alloy with a gradient composition structure lacks from the surface layer to the inside, the concentration difference of the bonding layer is relatively small, and the increase rate of the average particle size of the hard phase is not large, so the fracture toughness of the surface layer This is not only practical, but rather has the disadvantage of forming a nest inside the tissue, which is not practical.
Therefore, the present invention aims to increase the hardness and toughness of the surface layer portion even for products with complex shapes, and to achieve a composite structure with high internal strength. It is found that the particle size gradient of the hard particles and the concentration gradient of the bonding layer can be controlled accurately by controlling the particle size gradient and the concentration gradient of the bonding layer separately, rather than separately, and provide a desired ultra-hard material. is there.

The present inventors have found that an ideal high toughness cemented carbide is a skeletal structure with a small amount of bonding metal consisting of coarse hard particles in the surface layer, and a particle containing a large amount of bonding metal consisting of fine hard particles inside. On the other hand, an ideal high-strength cemented carbide must be composed of a dispersed structure, while the surface layer part is a skeletal structure with a small amount of bonded metal consisting of ultrafine and fine hard particles, and the inside is fine hard In view of the fact that it is composed of a particle-dispersed structure with a large amount of binding metal composed of particles, the present invention has been accomplished as a result of extensive research.
That is, the first invention is an M ₁₂ C type to M ₃ C type double carbide (M is one or more of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W, Fe, Co Carbide treatment is performed on a WC-Co-based green compact with the main component of the surface layer part, and then liquid phase sintering is performed on the surface layer WC using the liquid crystal sintering temperature as an index. A method for producing a cemented carbide material characterized by adjusting an average particle size.
The present invention uses the same starting material, and finer the surface layer portion that has been sintered with the liquid phase sintering temperature as an index. A double carbide with a composition of M ₁₂ C to M ₃ C is formed and carburized to decompose the double carbide to produce extremely fine and active WC particles, so the liquid crystal sintering temperature in the final liquid phase sintering Using the index as an index, WC particles that are 0.3 to 0.7 times finer than the inner part to WC particles that are 1.5 to 10 times coarser than the inner part can be formed on the surface layer of the sintered body.

Furthermore, the present inventors coated boride and silicide on the surface layer portion of the sintered body for the purpose of improving the surface layer hardness and imparting compressive residual stress, and 1200 to 1350 ° C. below the liquid phase sintering temperature. It was found that a high strength cemented carbide with a surface layer portion having a low friction coefficient that is extremely tough due to the gradient of the concentration of the binder phase from the surface layer portion to the inside can be obtained by performing diffusion heat treatment in the temperature range of . Therefore, the second invention is an M ₁₂ C type to M ₃ C type double carbide (M is one or more of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W, Fe, Co , Ni represents one or more of Ni) on the surface of the sintered body obtained after liquid phase sintering of a WC-Co-based green compact with the main component of the surface layer, The present invention provides a method for producing a high-strength cemented carbide comprising applying a compound containing silicon and subjecting it to a diffusion heat treatment in a temperature range of 1200 to 1350 ° C. below the liquid phase sintering temperature. According to the second invention, M ₁₂ C type to M ₃ C type double carbide (M is one or more of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and an iron group metal) WC-Co-based sintered tool comprising at least one of Fe, Co, and Ni) as a main component of the surface layer portion, and containing boron B or silicon Si by 0.010 to 1.0% by weight A high-strength cemented carbide sintered material characterized in that it has a surface layer portion included in the range, and the surface layer portion has hard particles having a distribution density higher than that of the inner quality portion.

Further, the third invention combines the first invention and the second invention to provide a cemented carbide material having both a hard particle size gradient and a binder phase concentration gradient from the surface layer portion toward the inside. M ₁₂ C type to M ₃ C type double carbide (M is one or more of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and any of Fe, Co, Ni) WC-Co compacted body with the main component of the surface layer part being carburized and then liquid phase sintered on the surface of the sintered body, melting point lowering element A compound containing boron or silicon is applied, and diffusion heat treatment is again performed in a temperature range of 1200 to 1350 ° C. below the liquid phase sintering temperature. According to the third invention, M ₁₂ C type double carbide (M is one or more of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and any of Fe, Co, Ni) WC-Co-based sintered body having a main component of the surface layer part, the average WC grain size of the surface layer part being 0.3 to 0.7 times smaller than that of the inner part It has a gradient of the structure in which the bonding metal in the surface layer has moved to the inner side, the hardness of the surface layer is HRA = 91 to 95, and the toughness is K _IC = 15 to 23 MN / m ^3/2 A high-strength cemented carbide sintered tool having excellent mechanical properties can be obtained. M ₃ C type double carbide (M represents one or more of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and one or more of Fe, Co, Ni) Is a WC-Co-based sintered body with a surface layer portion as a main component, and the surface layer portion WC has an average grain size that is 1.5 times larger than that of the inner portion, and the bonding metal of the surface layer portion High-strength cemented carbide sintered tool with excellent mechanical properties such as concentration gradient moved to the side, surface hardness of HRA = 88 to 92, and toughness of K _IC = 20 to 30 MN / m ^3/2 Is obtained.

As described above, according to the present invention, it is possible to provide a sintered tool having a hybrid structure having completely different characteristics in the surface layer portion and the inner quality portion, and the hardness, wear resistance, and toughness of the resulting cemented carbide. Excellent in fracture resistance and heat crack resistance.
Further, according to the present invention, a high toughness cemented carbide having a processed surface formed of coarse hard particles is provided, and the processed surface is formed of fine hard particles for a cutting blade, a progressive die, and a drawing tool. A high-hardness cemented carbide can be provided. Other applications include cold / warm / hot forging tools, canning tools, rolls, mining tool bits, crushing blades, cutting blades, and other wear-resistant tools.

(First embodiment)
In the present invention, M ₁₂ C to M ₃ C type double carbide (M is one of at least one of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and any of Fe, Co, Ni) The present invention can be widely applied to a WC-Co-based sintered body whose main component is a surface layer portion, but the following embodiment will focus on the WC-Co sintered body.
First, WC powder, Co powder, and other additive powders are milled to obtain a uniformly dispersed mixed powder, to which Wax as a lubricant is added to prepare a raw material.
Next, this raw material is compacted into a predetermined size and shape, pre-sintered for the purpose of removing Wax, and further shaped into an additional size and shape to complete a near-net-shaped molded product. This molded product has a porosity of 30-50 vol%.
As the next step, a double carbide phase having the following phase form is formed on the surface layer portion of the molded product within a depth of 3 to 5 mm from the surface at a volume ratio of 50 vol% or more.
M ₁₂ C [Co ₆ W ₆ C], M ₆ C [Co ₃ W ₃ C, Co ₂ W ₄ C], M ₃ C [Co ₃ W ₉ C ₄ ]
(The Co element may be replaced with the Fe and Ni elements, and W may be a solid solution with Ti and Ta.)

There are various methods for forming this double carbide. For example, by oxidizing the surface layer with various acids and then heat-treating it, a self-reducing reaction occurs to form a double carbide phase, or W ions are adsorbed on the surface layer using a W salt solution, and then A double carbide is similarly formed by heat treatment, and there is also a means for forming a double carbide by a method in which the carbide is deposited on the surface layer as a chloride and heat-treated. Aside from these means, in short, the composition of the surface layer portion may be put in the WC-γ-η three-phase region in the Co-WC ternary phase diagram. Here, the formation of the M ₁₂ C double carbide phase is necessary for the refinement of the surface layer particles of the final sintered body, and the formation of the M ₃ C double carbide phase is necessary for the grain coarsening.

  Then, carburizing heat treatment is performed to decompose the double carbide phase to form a fine and active WC phase. This is because by supplying carbon (C) to the double carbide phase in the temperature range of 600 to 1100 ° C., the double carbide phase decomposes and changes to a WC + Co 2 phase, so that ultrafine WC particles can be obtained.

Here, the carburizing treatment of the M ₁₂ C double carbide phase needs to be performed on the low temperature side, and the carburizing treatment of the M ₃ C double carbide phase needs to be performed on the high temperature side.
Alternatively, a nitriding heat treatment can be performed at this stage. Nitriding of ordinary WC particles is extremely difficult, but the nitriding reaction of fine and active WC particles generated by decomposition of the double carbide phase is considered to be almost equivalent to carburizing, and in the same temperature range, WCN other than WC + Co , WN can be generated easily.

Finally, liquid phase sintering is performed at a temperature of 1300-1500 ° C. to control the particle size of the surface layer WC particles. The WC particles are refined by low-temperature sintering at 1350 ° C, and the coarse particles are sintered at a high temperature range of 1400 ° C or higher. In 1350 ° C low-temperature sintering, the fine and active WC phase crystallizes and forms new nuclei, which increases the number of crystal growth nuclei together with the undissolved WC particles in the parent phase. As a result, a fine WC phase smaller than the fine WC grains in the inner part is generated in the surface layer part.
On the other hand, in high-temperature sintering at 1400 ° C or higher, the very fine and active WC phase preferentially dissolves during liquid phase sintering and preferentially precipitates on larger existing WC particles to grow.
The degree of grain growth is affected by the double carbide composition, and the higher the bond carbon content ratio, the greater the tendency of grain growth.
[Growth tendency] M ₁₂ C <M ₆ C <M ₃ C

The composite material obtained in this way has a particle diameter control region depth of 0.5 to 4.5 mm in the surface layer part, and the particle size is 0.3 to 0.7 times the internal particle size with fine particles and 1.5 to 10 with coarse particles. Double the size.
Further, the amount of the bonded metal at this time is a metallurgical action to keep the WC interparticle distance constant, so the difference in hardness between the surface layer portion and the internal portion where the particle diameter is controlled hardly changes.

  As an additional step, a boron compound or silicon compound powder is applied to the surface of the obtained sintered body and subjected to diffusion heat treatment in a temperature range of 1200 to 1350 ° C., so that the bonding metal in the surface layer is bonded to boron or silicon. It reacts to form a liquid phase, and boron and silicon diffuse into the solid phase region at the interface between the solid-phase bonded metal and the liquid phase, so that the liquid phase in the solid phase region progresses and the liquid phase moves inside. For this reason, the surface layer portion has an extremely small amount of bound metal, and a metal-rich structure can be obtained inside.

The final properties are that the surface layer hardness is HRA = 88 to 95, the toughness is K _IC = 15 to 30 MN / m ^3/2 , the mechanical properties of high hardness and high toughness are given, and the inside is high strength Mechanical properties are imparted. Furthermore, since the compressive residual stress acts on the surface layer region, it is optimal for various forging tools, press tools, mining tools related applications with high surface load stress.

    In the following, description will be given by taking, as an example, a die for a helical gear of a cold forging die and an excavating tool cutter bit.

[Die prototype for helical gear]
As shown in FIG. 1, the helical gear has a helical shape with a gentle screw part, and a pinion shaft of an automobile is typical as a product application. Conventionally, it has been manufactured by cutting, but in recent years it has been manufactured by cold forging. However, because forging is performed at an extremely high pressure, seizure and cracks occur in the mold teeth at an early stage, and the life is extremely short. In order to solve this problem, the alloy of the present invention is applied. .

1) Prototype production
Prepare 30 kg of metered raw material with a standard composition of WC-15% Co, adjusted to C / WC = 4.0% using 1.5μWC powder and 1.1μCo powder, and perform 30 Hr attritor milling using an alcohol solvent. Thereafter, paraffin wax is kneaded and granulated and sieved to obtain a finished powder.

Press molding In order to obtain the final sintered material size φ55 × 115L, press molding aiming at linear shrinkage F = 1.25 is made, and a compacted body of φ75 × 170L is produced.

Primary pre-sintering The dewaxing conditions were performed in a temperature range of 350 to 400 ° C. under an N2 carrier gas atmosphere, and the pre-sintering was performed under heat treatment conditions of 850 to 900 ° C. × 2 Hr in a vacuum atmosphere. Note that no shrinkage behavior occurs under these temperature conditions.

Forming process The shrinkage ratio of the pre-sintered body was accurately calculated to calculate the processing dimensions, and the NC was turned into a shape approximately 1.25 times larger than the sintered material dimensions shown in the diagram. In addition, about the shape of a blade part of an internal diameter, it did not perform a shaping | molding process, but only the cylindrical shape process.

Secondary pre-sintering Here, pre-sintering in a vacuum atmosphere of 1100 ° C. × 1 Hr was performed in order to improve the strength of the compact.

Immersion treatment
A 40% aqueous solution of tungstic acid (H ₂ WO ₄ ) was used as a combination of W supply and oxidant supply. The procedure is to fill the impregnating solution until the compact is sufficiently immersed in a stainless steel tray large enough to contain the compact, and soak the compact for 30 seconds. After the infiltration treatment, the removed molded body is immediately dried in a dryer at a temperature of 120 ° C.

Reduction heat treatment In this example, heat treatment was performed at 1000 ° C. × 2 Hr in a vacuum atmosphere. As a result of X-ray diffraction by TP, in addition to WC and Co phases, two-phase double carbides of Co ₆ W ₆ C [M ₁₂ C] and Co ₃ W ₃ C [M ₆ C] are confirmed in the surface layer. It was done.
In order to obtain a fine grained structure in the final liquid phase sintering, the presence of M ₁₂ C type double carbide phase is indispensable. For this purpose, the reduction heat treatment temperature is in the temperature range of 900 to 1100 ° C. Is preferred.

Carburizing heat treatment By supplying a carburizing gas into the furnace in a specified temperature range, the double carbide phase generated in the impregnation region is decomposed to generate extremely fine WC and Co phases.
A preferable carburizing temperature range is 600 to 900 ° C., but the carburizing atmosphere conditions in this example were a temperature of 900 ° C. × 30 min and a CO + H ₂ gas flow rate of 20 ml / min. The gas to be used may be a carburizing gas, and the temperature range is the solid phase region of WC-Co. Therefore, the phase transformation from double carbide to WC + Co is extremely stable and easy.
However, if the treatment temperature is raised to 1100 ° C or higher, the solid solution of carbon in the Co phase proceeds, so the possibility of generating free carbon in the alloy structure in the subsequent liquid phase sintering increases.

[Nitriding process]
In the above process, a nitriding heat treatment can also be performed. By subjecting the resulting double carbide phase to N ₂ , N ₂ + NH ₃ gas nitriding, extremely fine WCN and WN phases can be generated in addition to the WC and Co phases due to decomposition of the double carbide phase.
As the nitriding atmosphere conditions, a temperature of 800 to 1000 ° C. × 1 to 3 Hr and a gas flow rate of about 20 to 100 ml / min are preferable, and in the subsequent liquid phase sintering, in order to prevent N ₂ degassing from the material, it is usually used. What is necessary is just to maintain the partial pressure in the furnace below the pressure. As a result, the growing particles have a cored structure in which the inside is WC and the growing part is WCN or WN, and it has extremely excellent heat resistance.

Liquid-phase sintering It was processed in a vacuum sintering furnace at a temperature of 1350 ° C x 1.5Hr. In low temperature sintering at 1350 ° C, the fine and active WC phase crystallizes and forms new nuclei, which increases the nucleus of crystal growth along with the undissolved WC particles of the parent phase. As a result, a fine WC phase smaller than the fine WC grains in the inner part is generated in the surface layer part. As a result of tissue observation, a microstructure refined to 0.5 to 1.0 μm was confirmed in the surface layer region including the inner surface.

Application of boron compound Apply an alcohol slurry of 20% BN concentration to the inner surface of the sintered material obtained in this way, and dry it for 1 hour in a dryer set at a temperature of 40 ° C.

Diffusion heat treatment After coating and drying, the material is subjected to diffusion heat treatment of 1300 ℃ × 2Hr. Since a boride concentration gradient is formed from the surface to the inside, the liquid phase in the surface layer continues to diffuse into the interior, and finally there is almost no bound metal remaining in the surface layer region, and there is a metal rich in the inside. A simple structure is formed.
The mechanical properties of the developed alloy thus obtained are as follows when the surface layer and the interior are roughly divided.

Table 1

Preparation of comparative alloy As a comparison with the newly developed alloy, a cemented carbide material of the same size and shape was prepared using a WC-11% Co alloy based on 1.5μWC. The procedure is to prepare a WC-11% Co mixed raw material, and after press molding
After pre-sintering at 900 ° C. and forming into a desired shape, 1380 ° C. × 1 Hr vacuum sintering was performed to produce a material.

Mold processing into helical gear shape The mold shown in Fig. 2 was manufactured. The casing material for protecting the newly developed cemented carbide was 8 types of SNCM, and the allowance for cemented carbide was 0.5%. The inner diameter of the cemented carbide was processed into a helical gear shape by electrical discharge machining using a Cu-W electrode molded into a male shape, and a final finishing lapping with a third grade accuracy was performed.
After finishing the inner surface of the alloy, it is removed from the casing, TiC + TiN CVD coating is applied, the casing is re-executed, and the finished mold is finished.

Actual machine evaluation All conventional die dies are coated with CVD (TiC + TiN), but here we also compared CVD treated and untreated products.
The results are shown in the table below. In the comparative alloy, the CVD-untreated dies produced extremely early seizure and had the shortest life, and the longest-life dies were developed alloys that were untreated with CVD.
The reason why the lifetime of the CVD-treated developed alloy did not extend due to the missing teeth was thought to be that the coating film cracked and propagated to the cemented carbide base material.
Therefore, this developed alloy is an ideal tool material with excellent wear resistance without any coating treatment, tough structural characteristics, excellent fracture resistance, and dramatically improved fatigue life. It became clear that there was.

Table 2

[Prototype casing bit]
The casing bit is a bit used for foundation work for a building structure. As shown in Fig. 3, it is a drilling tool in which a cemented carbide is brazed to an S55C support metal. This tool is attached to the tip of a steel pipe, applies a load while rotating the pipe, and excavates the ground from the ground toward the ground. The depth of excavation is the depth to reach a bedrock layer with sufficient strength. For example, when the depth is up to 30m, the excavation is carried out by connecting steel pipes. The drilling performance is largely governed by the characteristics of the cemented carbide brazed to the bit, and conventionally, coarse-grained cemented carbide has been mainly used to avoid the fracture of the cemented carbide. However, since the excavation is performed at an extremely high pressure, wear of the cemented carbide blade portion progresses at an early stage, which hinders maintenance of excavation capability. On the other hand, when medium to fine grain cemented carbides are used, chipping and destruction of the cemented carbide blades often proceed rapidly, and in this case, excavation does not proceed at all and the major problem is that the construction period is delayed. Had occurred. In order to solve these problems, the alloy of the present invention was applied. The expected mechanical properties were set to HRA = 90-91.5 for the surface layer hardness and KIC = 20-25 MN / m ^3/2 for fracture toughness.

Prototype of raw material Here, the raw material used for the trial manufacture of the die for the helical gear was used.

Press molding In order to obtain the final sintered material size of 40 × 22 × 40, press molding aiming at a linear shrinkage of F = 1.25, and a compact of 50 × 100 × 150 is produced.

Primary pre-sintering The dewaxing conditions were performed in a temperature range of 350 to 400 ° C. in an N 2 carrier gas atmosphere, and the pre-sintering was performed in a vacuum atmosphere at 850 to 900 ° C. × 2 Hr.

Forming processing The processing dimensions were calculated from the shrinkage rate of the pre-sintered body, and were processed into a shape approximately 1.25 times the size of the sintered material using various cutting machines and grinding machines using diamond tools.

Secondary pre-sintering Here, in order to improve the strength of the compact, vacuum atmosphere pre-sintering of 1100 ° C. × 1 Hr was performed.

Immersion treatment A 30% aqueous solution of ammonium metatungstate (AMT) and cobalt nitrate was used here. The immersion time of the molded body was 20 seconds. After the immersion treatment, the removed molded body is immediately dried in a dryer at a temperature of 120 ° C.

Reducing heat treatment Heat treatment was performed at 1300 ° C. × 1 Hr in a vacuum atmosphere. As a result of X-ray diffraction by TP, two-phase double carbides of Co ₂ W ₄ C [M ₆ C] and Co ₃ W ₉ C ₄ [M ₃ C] in addition to the WC and Co phases in the surface layer region Was confirmed. However, in the temperature range of 1300 ° C. or higher, the compacting of the molded body proceeds, and therefore, the progress of carbon internal diffusion becomes extremely slow during the subsequent carburizing treatment.

Carburizing heat treatment Carburizing atmosphere conditions were as follows: temperature 1100 ° C x 30 min, CO + H ₂ gas flow rate 20 ml / min. The gas used may be a carburizing gas, and the temperature range is the solid phase region of WC-Co. Therefore, the phase transformation from double carbide to WC + Co is extremely stable and easy.

Liquid-phase sintering It was processed in a vacuum sintering furnace under the temperature condition of 1420 ° C x 1Hr.

Application of Boron Compound On the outer surface of the sintered body material obtained in this way, an alcohol slurry with a B ₄ C concentration of 20% was applied and dried for 1 hour in a dryer set at a temperature of 40 ° C.

Diffusion heat treatment After coating and drying, the material was subjected to diffusion heat treatment at 1300 ° C x 2 Hr. Eventually, almost no binding metal remains in the surface layer region, and a metal-rich structure is formed inside.
The mechanical properties of the developed alloy thus obtained are as follows when the surface layer and the interior are roughly divided.
In addition, as a comparative alloy, a bit sample and TP were prepared and compared using a WC-14% Co alloy with a WC particle size of 6 μm.

Table 3

Casing bit production
S55C A support metal made by cutting from a forged product was heat-treated and adjusted to HRC = 35-40 hardness, and then a cemented carbide material was high-frequency brazed into a blade shape to complete a casing bit. There are L and R types of bits. The shape shown in the schematic diagram is the R type. The arrangement for attaching the bit to the end of the pipe is generally in the order of -RRLRRL-, and the attachment to the casing pipe was performed in this order.

Evaluation of actual machine The casing pipe used for excavation is 2200mm in diameter, and a total of 36 bits are used at the tip. The breakdown was 24 R type and 12 L type. According to the results of the geological survey, there were gravel layers and boulders from 8m to 12m deep, and the average excavation depth of the foundation pile was about 18m. The life evaluation of the bit was performed by the number of exchange bits per foundation pile. In other words, when the excavation of the 18m foundation pile was completed, the entire pipe was taken out, the worn state of the bit was confirmed, and those that were deemed necessary for replacement were replaced.

These results are shown in the table below. Obviously, the life of the developed alloy bit is 11 to 18 times longer, which is a longer life than the comparative material.
Table 4

(Second Embodiment)
A sintered tool is integrally formed from an inner part and a surface layer part formed by heat treatment so as to surround the inner part. Basically, the inner part consists of a hard particle and a binder metal that binds these particles. In the second embodiment, the surface layer portion necessarily includes hard particles and boron B and / or silicon Si. The surface layer portion may contain a binder metal, but it is preferable that the content is less than or substantially not contained in the inner quality portion in order to increase the surface hardness.

Hard particles in the sintered tool include carbide, nitride, or carbonitride, and in particular, WC, TiC, TaC, NbC, VC, Cr ₂ C ₃ as a carbide, TiN, TaN, NbN as a nitride, At least one or more of VN, Cr ₂ N, and ZrN are used.

The other binder metal is at least one selected from iron group metals, that is, Fe, Ni, and Co. From the viewpoint of corrosion resistance, heat resistance, and oxidation resistance, Ni or Co can be preferably used. Ni and Co form a solid solution of B in the surface layer portion, and form hard borides NiWB and CoWB in the presence of WC, thereby contributing to surface hardening. In the case of silicon Si, Ni and Co solid-dissolve Si in the surface layer part, and form hard silicides NiWSi ₄ and CoWSi ₄ in the coexistence of WC, thereby contributing to surface hardening.

  About an internal part, it is a hard particle, a binder metal, and a sintered compact, The ratio of content of a binder metal and a hard particle exists in the range from 5:95 to 40:60. When the content ratio of the hard particles is lower than 5:95, the binder metal is too small to form a sintered body. When this content ratio is larger than 40:60, there are few hard metals and a sintered compact cannot fully be hardened.

  The content ratio of binder metal to hard particles is preferably in the range of 5:95 to 30:70. The ratio of this content is selected depending on the use of the sintered tool, but in general, in applications that require toughness, particularly impact resistance, together with surface hardness, Among them, the hard particles are reduced and the content ratio of the binder metal is high and prepared. On the other hand, for applications that particularly require surface hardness and wear resistance, the hard particle content ratio is increased within the above-mentioned content range.

  On the other hand, as will be described later, the surface layer portion of the sintered tool is formed by diffusing boron B and / or silicon Si from the surface of the sintered body in the heat treatment process of the sintered body having the above composition. A Si-containing layer is utilized.

  In the present invention, the surface layer portion contains boron B or silicon Si alone or in a total weight and is in the range of 0.010 to 2.0%, and the surface layer portion is harder than the internal portion. The distribution density is high. In particular, the boron or silicon content in the surface layer is preferably in the range of 0.050 to 1.0%. When both boron and silicon are included, the total amount is preferably within the above range.

  Binder metal is reduced from the inner part. The content of boron B or silicon Si is set to 0.010 to 2.00% in order to ensure the hardness of the surface layer portion. When the content of boron or silicon is less than 0.010%, the surface layer portion is subjected to the diffusion heat treatment. The diffusion of the binder metal from the inside to the inside becomes insufficient. On the other hand, when it exceeds 2.00%, the surface layer portion cannot follow the volume change accompanying the internal diffusion of the binder metal phase, and surface cracks occur during the diffusion heat treatment. This is because it tends to occur. By setting the boron or silicon content to 0.050 to 1.0%, it is possible to increase the diffusion of the binder metal from the surface layer portion to the inside, and to effectively prevent surface cracks and the like. As a result, the surface layer portion has a relatively small binder metal content and a high hard particle content as compared with the internal portion. This makes it possible to reduce the average interval between adjacent hard particles, which is also estimated by volume, where the distribution density of the hard particles is higher than that of the inner part, and due to the high density of hard particles. The surface hardness is higher than that of the inner part.

  The distribution density of the hard particles is the highest near the surface in the surface layer portion, is reduced in the depth direction of the surface layer portion, and approaches the distribution of the inner quality portion. Along with the gradient distribution of such hard particles, the content of the binder metal is made lower in the surface layer portion than the inner portion, and the hardness distribution is also inclined so as to decrease from the surface vicinity toward the inner portion. ing.

  The binder metal element content is an average value in a range from the surface of the surface layer portion to a depth of 0.5 mm, and is preferably 2% or less by weight. In this way, the surface layer portion of the tool of the present invention substantially consists of a hard particle phase and a boride and / or silicide phase, and the tool is hardened by agglomeration of the hard particles and boron and / or silicon compound. High surface hardness can be obtained on the surface.

  In the sintered tool of the present invention, the average particle size of the hard particles in the sintered tool is preferably in the range of 0.2 to 15 μm. As the hard particles are made finer, the hardness increases. However, if the particle size is smaller than 0.2 μm, the amount of change in the bonded carbon and nitrogen in the hard particle phase increases, and the stability in terms of surface hardness cannot be maintained. . On the other hand, if it exceeds 15 μm, it is preferable to avoid it because the wear resistance is lowered. Although the particle diameter of the surface layer portion and the inner quality portion varies depending on the application and shape of the tool, in particular, the average particle size is more preferably in the range of 0.5 to 10 μm.

  As described above, in the surface layer portion, the binder metal content is reduced, and in the surface layer portion, fine hard particles are densely distributed, and the surface layer portions are adjacent to each other from the inner quality portion. The average interval between hard particles can be reduced. Such a surface layer microstructure is useful for increasing the hardness of the surface layer portion composed of hard particles including borides, reducing the friction coefficient, and increasing the wear resistance and heat resistance strength.

In this surface layer portion, as described above, boron is contained together with the hard particles, but boron combines with the binder metal to form an iron group metal boride, and the boride exists as a precipitated phase between the hard particles. Thus, the iron group boride itself is hard, and therefore, hardening due to the contribution of the iron group boride is recognized in the surface layer portion. The boride includes FeWB, NiWB, or CoWB in the presence of WC. The silicide includes NiWSi ₄ and CoWSi _{4 in} the presence of WC.

As described above, the sintering tool can use hard particles, WC as a main component or TiC or a mixture thereof, and can use Ni or Co as a binder metal. As an example of the tool, when the hard particles are WC and the binder metal is Co, the inner part is composed of a WC phase of a fine particle phase as a main phase and a metal Co phase (Co solid solution) in a required blending amount. The surface layer portion includes a WC phase and a finely precipitated CoWB phase as a boride phase (if the Co phase is present, a very small amount of Co solid solution phase). . Further, as a silicide phase, a CoSi ₂ phase, a WSi ₂ layer, and a CoWSi ₄ layer that are finely precipitated are included in the surface layer portion.

  The surface hardness of the WC-Co-based sintered tool of the present invention depends on the hardness of the internal part, but is particularly Hv 1000 or more, usually in the range of Hv 1400 to 1800, or more, for example, Hv 2300. Those having the following are preferred.

  In general, the thickness of the surface layer portion is 2 mm or more, preferably the distance from the surface to the position where the linear portion of the hardness distribution curve from the surface reaches the average hardness of the inner quality portion is 2 mm or more. Secures 4 mm or more.

  In this way, the surface layer portion of the present invention achieves surface hardening by increasing the density of hard particles and the coexistence of the iron group metal boride, and the inner portion is required by the required blending of the hard particles and the binder metal. The toughness, hardness and strength can be ensured.

  Regarding the method for producing a sintered tool of the present invention, first, a sintered body is produced. The sintered body is a green compact of a desired shape by compression molding a mixed powder of hard particles and an iron group binder metal. Then, the green compact is made into a normal sintered body by usual liquid phase sintering. Thereby, a densified and uniform sintered body is obtained. This sintering method is entirely sintered using a conventional method. After sintering, the sintered body can be appropriately machined such as precisely cutting, grinding, electric discharge machining into a desired shape.

  Next, a boron or silicon coating layer is formed on the surface of the sintered body. In order to form this kind of coating layer, a boron coating containing boron is coated, and in the heat treatment, the sintered body having the boron coating layer is heated to form a surface layer portion rich in boron or silicon.

  In this heat treatment, the sintered body having the boron coating layer is lower in the vacuum or in an inert gas, preferably in a nitrogen gas atmosphere, at a temperature lower than the liquid phase temperature in the sintered body, and the sintered body. Heat and hold for a desired time in a temperature range higher than the eutectic temperature of the boron-containing phase therein. During the heat treatment, boron in the boron coating layer is diffused from the surface of the sintered body to the inside to form a surface layer portion rich in boron, and the melt in the surface layer portion is diffused and transferred to the internal material portion, and the sintered body The distribution density of the hard particles in the surface layer portion is made higher than that in the inner quality portion, and after cooling, boron or silicon is precipitated in the surface layer portion as a boride and / or silicide phase containing a binder metal and hardened. A sintered tool having a part is obtained.

Regarding the details of the manufacturing method of the sintered tool of the present invention, as described for the sintered tool, the hard particles include carbide, nitride, or carbonitride, and in particular, as carbide, WC, TiC, TaC, NbC, VC, Cr ₂ C ₃ and at least one or more of TiN, TaN, NbN, VN, Cr ₂ N, and ZrN are used as nitrides. The other binder metal is at least one selected from iron group metals, that is, Fe, Ni, and Co. Preferably, Ni and Co can be used.

  When Ni or Co as the binder metal contains B or Si, Ni—B or Ni—Si alloy or Co—B or Co—Si alloy alloy, or Ni—W—B or Ni—W—Si alloy or Co— Since the eutectic temperature of the WB or Co—W—Si alloy is lower than the alloy solidus temperature of Ni or Co and the above carbide, the Ni—WB or Ni—W—Si alloy or Co -W-B or Co-W-Si alloy is used for heat treatment, and as described later, the distribution of hard particles in the surface layer portion is made higher than that in the inner portion, and is used for surface hardening.

  The ratio of the content of the hard particles and the binder metal in the hard particle raw material and the binder metal raw material powder is preferably in the range of 5:95 to 30:70. The ratio of this content is selected depending on the use of the sintered tool, but in general, in applications that require toughness, particularly impact resistance, together with surface hardness, Among them, the hard particles are reduced and the content ratio of the binder metal is high and prepared. On the other hand, for applications that particularly require surface hardness and wear resistance, the hard particle content ratio is increased within the above-mentioned content range.

  The hard particles of the raw material have an average particle diameter of preferably 0.2 to 15 μm, and preferably 0.5 to 10 μm.

  By using the above raw material hard particles, the particle size of the surface layer part and the internal part in the product tool can be obtained by sintering and heat treatment, although it varies depending on the application and shape of the tool. The average particle diameter of the particles is an average particle diameter, and a range of 0.2 to 15 μm is used. As described above, as the hard particles are made finer, the surface hardness increases. However, if the particle size is smaller than 0.2 μm, the amount of change in the bonded carbon and nitrogen in the hard particle phase increases, and the surface hardness is stable. Sex cannot be maintained. On the other hand, if it exceeds 15 μm, the wear resistance is lowered, so it should be avoided. Although the particle diameter of the surface layer portion and the inner quality portion varies depending on the application and shape of the tool, in particular, the average particle size is more preferably in the range of 0.5 to 10 μm.

  The mixed powder of hard particles and binder metal is compression-molded into a green compact having a desired shape, and the green compact is sintered in the same manner as a conventional sintered part. Sintering is performed after pre-sintering and then main sintering is performed to obtain a dense sintered body. For example, conventional liquid phase sintering can be applied.

In the boron or silicon coating step of the present invention, a coating agent containing boron or silicon is applied to the surface of the sintered body, and the boron coating material for this purpose contains a boron compound, and boron oxide, nitride or Contains carbides or their precursors, such as carbonates and hydroxides. For example, SiB ₆ , BN, B ₄ C, B ₂ O ₃ , H ₃ BO ₃ , borane, or an organic boron compound can be used for the coating material. The silicon coating material includes a silicon compound, and includes a carbide or nitride, a boride, or a precursor thereof, or an intermetallic compound. More specifically, Si, SiH4, SiCl4, SiC, Si3N4, SiB6, or CoSi2, MoSi2, CrSi2, WSi2, or silanes, polysilane polymers, and other organosilicon compounds may be mentioned.

  The boron coating material contains these boron compounds and coats the surface of the sintered body. The coating material may be applied directly to this surface, but from the certainty of coating, these boron compounds are preferable. Is suspended in water or a non-aqueous solvent to prepare a slurry-like coating solution, which is applied to the surface of the sintered body. The application may be performed by, for example, brushing the coating liquid on the surface of the sintered body, spraying it with a spray or the like, or dipping the sintered body in a coating bath and pulling it up. Next, the coating liquid is dried on the surface of the sintered body to leave the coating material.

  The coating liquid may be applied to the entire surface of the sintered body, but the surface to be cured of the sintered tool is limited, and appropriate masking is applied to other surface portions to provide a boride-containing coating material. If the coating is prevented, the surface layer portion is formed only in a desired surface area by the heat treatment process, and the surface of the tool can be hardened by the surface layer portion. It is soft and can retain high toughness.

  On the other hand, as a boride or silicide coating process as another means, chloride, fluoride, hydride or organometallic compound is introduced into a heating furnace and decomposed, and vapor deposition coating is performed on the surface of the sintered body. There is also a method. This method is generally called a chemical vapor deposition method [CVD], but in addition to the conventional atmospheric pressure CVD method and the low pressure CVD method, recently, a plasma CVD method, a thermal CVD method, or a laser CVD method has been developed. The film formation rate by vapor deposition is improved to 0.1 lμm / sec or more.

  Materials used as raw material sources at this time include boron trichloride and silicon tetrachloride as chloride, boron trifluoride and silicon tetrafluoride as fluoride, and hydride is boron hydride. Examples of (borane) include diborane, pentaborane, dihydroborane, and derivatives thereof, and examples of silicon hydride (silane) include monosilane and disilane. Examples of organometallic compounds include organoboron compounds and organosilicon compounds, such as trialkylboron, chlorosilane, and alkoxysilane, and more specifically, trimethylboron, triethylboron, tri-n-propylboron, and tri- n-Butylboron and the like, and dichloromethylsilane, chlorodimethylsilane, chlorotrimethylsilane, tetramethylsilane and the like. Other compounds include organic boronic acids.

Specifically, these compounds are made into a gaseous state, and the gaseous compound is introduced into a heating furnace set at a furnace temperature at which the compound can be decomposed with a carrier gas at a predetermined flow rate, and is introduced into the surface of the sintered body. Borides or silicides from the decomposition of compounds are deposited. As the continuous decomposition / deposition reaction proceeds for a predetermined time, a coated metal layer having a predetermined film thickness is formed on the surface of the sintered body.
Adjustment of the film thickness at this time is controlled by gas concentration, carrier gas flow rate, heating temperature, heating time, and the like.

On the other hand, as another coating means, a dense boride or silicide metal coating can be formed by thermally spraying a powder aggregate of boride or silicide heated to a semi-molten state on the surface of the sintered body at high speed. . Examples of these borides and silicides include SiB ₆ , SiC, Si ₃ N ₄ , BN, and B ₄ C.

  In the heat treatment, the sintered body whose surface contains boron or silicon and is coated with the dry coating material is then heated while being held in vacuum. The temperature of the heat treatment should be lower than the solidus temperature or eutectic temperature determined from the alloy composition of the hard particles and the iron group binder metal, and the sintered body composition should The temperature is selected so as not to form a melt and higher than the eutectic temperature of the alloy system containing boron or silicon, hard particles, and a binder metal from the coating layer on the surface.

  That is, the present invention utilizes the fact that the eutectic temperature containing boron or silicon is lower than the eutectic temperature of the sintered body not containing boron or silicon, and the heat treatment temperature is set to a temperature between these eutectic temperatures. By setting, a part of the melt is formed only on the surface or the surface layer. This melt is composed of most of boron and iron group metals and a very small portion of hard particles, and most of the hard particles remain solid.

In the WC-Co-based sintered tool, from the phase diagram of the WC-Co pseudo-binary alloy, the eutectic temperature is about 1320 ° C, while the Co-B system has a Co-side eutectic point (ie, Since the eutectic temperature of Co—Co ₃ B) is about 1110 ° C., the heat treatment temperature is 1150 to 1310 ° C., preferably 1200 to 1300 ° C. is used.

Further, in the WC-Ni based sintered tool, from the phase diagram of the WC-Ni pseudo binary alloy, the eutectic temperature is about 1390 ° C, while the Ni-B based is the Ni side eutectic point (ie , Ni—Ni ₃ B eutectic temperature) is about 1090 ° C., and therefore, the heat treatment temperature is in the range of 1150 to 1380 ° C., preferably 1200 to 1370 ° C. between the two eutectic temperatures. A range of is used.

Furthermore, since both the TiC-Co system and the TiC-Ni system have a liquidus appearance temperature of about 1270 ° C., in the TiC-Co system and the TiC-Ni system sintering tool, the heat treatment temperature is 1200-200. 1250 ° C is preferred. Furthermore, since the eutectic temperature of the Mo ₂ C—Ni system is about 1250 ° C., the TiC—Mo ₂ C—Ni system diffusion heat treatment can also be carried out in the temperature range of 1200 to 1250 ° C. In the system, mixing of Mo ₂ C can suppress the growth of carbide grains and improve the sinterability in the TiC—Co or TiC—Ni system. The appearance of the liquid phase, the formation of compounds, and the diffusion transfer in the heat treatment process as described above are the same for silicon. The Co-side liquid phase appearance temperature of the Co-Si system is around 1200 ° C. The liquid phase appearance temperature drops to 1000 ° C or less with Ni-30% Si composition.
For these reasons, the silicon diffusion heat treatment temperature in the WC-Co alloy is 1250 to 1320 ° C, and the WC-Ni alloy is in the range of 1150 to 1350 ° C.

  When heat treatment is performed in the above temperature range, at the initial stage of the heat treatment, boron in the boron-containing coating layer coated on the surface of the sintered body reacts with the iron group metal on the surface, and the surface contains boron at a low temperature. However, since the inside of the sintered body does not contain boron, the melt remains a solid that does not melt at the processing temperature. As the heat treatment time elapses, the melt on the surface portion dissolves the internal metal and penetrates into the interior while accompanying boron. Along with the permeation and diffusion into the melt, the melt is reduced near the surface, and the concentration or distribution density of the hard particles is increased.

  The region where the boron or silicon content is high and the hard particle density is high is the surface layer portion, but the surface layer portion has a small interval between adjacent particles, and the remaining boron or silicon content. Also gets higher. If the substrate is cooled or allowed to cool after the desired treatment time, the surface layer portion forms a compound of boron or silicon and a binder metal, and a boride or silicide precipitates. The surface layer portion constitutes a layer composed of boride or silicide and hard particles having a high distribution density. However, in this manufacturing method, the hard particles in the surface layer portion are hardly grown and are densified. Curing of the surface can be realized.

The boron or silicon content in the surface layer portion after the heat treatment can be controlled by the type of boron or silicon compound in the coating material before the heat treatment and the boron or silicon coating amount per surface area of the sintered body. For example, boron in the boron coating layer is preferably in the range of 5.0 to 40 mg / cm ² with respect to the coating surface in terms of metal boron B element. In this range, the surface layer portion can contain boron B in the range of 0.050 to 0.50% by weight as described above. In the surface layer portion, such a high content of boron is because boron exists as a compound of an iron group metal. The same applies to silicon.

  When the manufacturing method of the present invention is applied to a WC-Co-based sintered tool, the surface hardness depends on the hardness of the inner part, but the Vickers hardness Hv 700 or more than the surface hardness of the inner part, In particular, Hv1000 or more, usually in the range of Hv1400 to 1800, or more, for example, those having Hv2300 are preferred.

  Generally, the thickness of the surface layer part is 3 mm or more, preferably the distance from the surface to the position where the straight line part of the hardness distribution curve from the surface reaches the average hardness of the internal part Can secure 6 mm or more.

  The sintered tool of the present invention can be widely applied to cutting tools, plastic working tools, rock drill bits for mining / civil engineering and the like.

  Examples of cutting tools include single tool blades, milling cutters, drills and reamers, but drills and reamers are sintered bodies of ultrafine particles with a hard particle size of 1.0 μm or less, and their tool length L Since the shape with a high ratio to the diameter D (L / D ratio) is required, a material with high toughness is required. However, the structure of the present invention has high toughness at the center and high hardness and microstructure in the surface layer. By doing so, the surface layer portion has a high hardness advantageous for the configuration of the cutting edge, and the tool life can be increased.

  Examples of the processing tool include a press die, a forging die, a punch, and the like, and the sintered tool of the present invention can be applied to them. As a mold, for example, a mold for can making has conventionally used a ceramic material or a Ni-based superalloy. However, a ceramic tends to cause surface defects, and a superalloy is difficult to prepare a metal structure. However, according to the present invention, the WC-Co-based sintered body is subjected to boron diffusion heat treatment to increase the distribution density of the hard particles containing boron, thereby increasing the hardness, high wear resistance, and adhesion resistance. Due to the corrosion resistance, a mold having a long mold life can be obtained.

  The processing tool also includes a drawing die for steel pipes and a plug for wire drawing, and conventional cemented carbide has a problem of seizure, and is used with a TiN coating on the surface of the cemented carbide to prevent seizure. In some cases, seizure is likely to occur, and by performing boron diffusion heat treatment using a WC-Co system as a sintering tool of the present invention, CoWB (or Si) of the surface layer portion reduces the friction coefficient, The adhesion resistance is improved and the tool life can be extended.

  Examples of other processing tools include hot extrusion dies for aluminum alloys, and the die is a sintered tool of the present invention, instead of the conventional hot mold steel, and the extrusion temperature is around 500 ° C. Thus, in the presence of the CoWB or CoWSi phase in the surface layer portion, the adhesion resistance is improved and the die life can be improved.

  Furthermore, the cold forging punch for backward extrusion has a large compression load and extremely high frictional force with the work material, and is used under severe conditions. For this reason, it is often used after coating treatment. However, the present invention can be applied here to prevent breakage accidents due to insufficient punch toughness, reduce seizure wear of the bearing portion of the punch, and improve tool life.

  Two types, Co containing 10% in WC and Co containing 20% in WC, obtained by mixing commercially available tungsten carbide WC powder having an average particle diameter of 1.5 μm and 1.3 μm metallic cobalt Co powder. To the mixture. The mixed powder is compression-molded, the green compact is subjected to intermediate sintering, and is molded so that the size after sintering is 30 mm diameter × 30 mm length, and then in vacuum at 1400 ° C. for 1 hour. Liquid phase sintering was performed to obtain respective sintered materials.

Next, boron carbide B ₄ C is used as a boron source for the heat treatment, and commercially available boron carbide B ₄ C is pulverized with ethanol for 30 hours using ethanol in order to prepare a boron-containing coating material. A slurry containing 9% ₄ C was prepared. Polyethyleneimine was added to the slurry to obtain a boron-containing coating solution for coating.

  A dipping method was used for the coating method, and the sintered material was taken out after being dipped in the coating solution, and then dried in a dryer at 40 ° C. to prepare a sample.

As a comparative example, the above-described sintered material was used as it was without applying a boron-containing coating material.
The above example samples and comparative example samples were subjected to diffusion heat treatment under the following conditions. The sample is held in a vacuum furnace, controlled to a furnace pressure of 40-80 Pa, heated at a heating rate of 5 ° C./min, and held at three heat treatment temperatures of 1200 ° C., 1250 ° C. and 1280 ° C. for 3 hours. Diffusion heat treatment was performed, and the furnace order was given later.

  The heat-treated sample was cut at a position of 15 mm in length, the cut surface was polished, the microscopic observation of the cross-sectional structure was performed, and then the hardness was measured with a Vickers hardness meter while changing the depth from the surface. It was.

Boron-coated WC-20% Co sintered tool was coated by dipping in 9% B ₄ C coating solution using fine hard particles (particle size of 1-2 μm) and diffusion by boron Regarding the cross-sectional structure of the heat-treated sample, as shown in FIG. 4 (A), in the structure photograph of the internal part, many clear white metallic Co phases are recognized in the WC particle group. FIG. 4B shows the structure of the surface layer portion of this sample, but it has a dense carbide WC and almost no white metal phase is observed. When these structures are compared, it is the result of the metal Co phase near the surface moving inside during the heat treatment process. Compared to FIG. 4 (A) and FIG. 4 (B), the surface layer part and the inside are In both cases, there is almost no difference in the particle diameter of the WC particles.

Similarly, a sintered body subjected to diffusion heat treatment with boron by using a coarse hard particle (particle diameter: 3 to 6 μm) having a WC-20% Co composition, dipped in a 9% coating solution of B ₄ C and coated. FIG. 5A shows a micrograph of the cross-sectional structure of the inner part and FIG. 5B shows the cross-sectional structure for comparison. 5 (B)) reduces the binder metal phase (appears as a white phase in FIG. 5 (A)) compared to the internal part (FIG. 5 (A)), but in both cases, hard particles (WC It can be seen that the particle diameter of the particles is almost unchanged.

  On the other hand, as for the structure of the comparative example which has not been coated, a large structural change was not recognized in both the surface layer part and the inside, similar to FIG.

  Next, the hardness measurement results are shown in Table 5 and FIG. As is clear from the figure, a clear gradient was recognized in the hardness distribution in the material of the coating treatment. In the above heat treatment range, it can be seen that the lower the temperature, the higher the surface hardness and the smaller the surface layer thickness. When the heat treatment temperature is increased, diffusion of the melt into the interior proceeds, the surface layer portion is relatively thick, and the surface hardness tends to decrease. That is, the difference in hardness between the surface layer portion and the inner quality portion is about HV = 300 to 600, and the sample having a higher heat treatment temperature has a larger gradient depth.

Table 5

  The hardness gradient region is also a boron B diffusion region, and it was considered that the internal diffusion of boron B progressed by increasing the heat treatment temperature. The main factor for improving the hardness of the surface layer is that the interparticle distance on the surface layer side is reduced due to the decrease in the metal phase of the surface layer, and the hardness improvement effect due to the formation of CoWB is also considered to contribute. It is done. As a matter of course, a substantially uniform hardness distribution was obtained for the untreated product.

  When a sample having a thickness of 2 mm was cut out from the surface layer portion and the boron B content was measured by an ICP-MS method, an analysis result of 280 to 330 mg / kg was obtained, and diffusion of B could be confirmed.

The sintered material prepared in Example 3 was used for coating under the coating conditions in which the B ₄ C slurry concentration was 9%, 18%, and 24%, and the heat treatment was performed at a heating rate of 5 ° C. / Min, and a heat treatment temperature of 1280 ° C. for 3 hours.

  After the obtained sample was cut and polished at the center, the cross-sectional structure was observed, and then the hardness was measured with a Vickers hardness meter while changing the depth from the surface. The results are shown in Table 6 and FIG.

Table 6

  Referring to Table 6 and FIG. 7, both WC-10% Co and WC-20% Co using tungsten carbide WC powder having a particle size of 1.5 μm are compared with Example 1 in terms of diffusion depth. It is as large as 2 to 5 mm, and it can be seen that the diffusion depth increases in proportion to the coating material concentration.

  Thus, it can be seen that the hardness distribution can be appropriately obtained in the surface layer portion by setting the coating material concentration, that is, the surface addition amount of boron, and the heat treatment temperature conditions.

  The sample heat-treated in Example 4 was subjected to X-ray diffraction of the surface layer portion. Although not shown, a diffraction peak corresponding to CoWB was observed in the diffraction chart. From this, it is considered that the effect of the hard boride particles contributed to the improvement of the hardness of the surface layer portion.

Next, a commercially available WC powder having an average particle diameter of 0.55 μm, a metal Co powder having an average particle diameter of 1.3 μm, a chromium carbide Cr ₃ C ₂ powder, and a vanadium carbide VC powder are mixed, A mixed powder having the composition WC-20% Co-0.7% Cr-0.4% V was made and compacted to obtain a compact. In the same manner as in Example 3, the green compact was subjected to intermediate sintering and then cut into a cylindrical body having a diameter of 30 mm and a length of 30 mm. Similarly, vacuum sintering at 1350 ° C. × 1 hour was performed for testing. The sintered material was used.

As the boron coating material, a slurry-like coating liquid containing boron carbide B ₄ C was used in the same manner as in Example 3. Further, commercially available hexagonal boron nitride (h-BN) was added to ethanol for 30 hours. Was then pulverized by ball milling, and polyethylene imine was added to the obtained 9% h-BN slurry to prepare a coating solution for BN coating.

  The sintered material was subjected to two types of coating: a BC-containing slurry coating process and a coating process with a BN-containing slurry coating solution. On the other hand, the sintered material of WC-10% Co and WC-20% Co prepared in Example 1 was subjected to BN coating treatment. After drying, all samples were subjected to diffusion heat treatment at 1280 ° C. for 3 hours.

  The heat-treated sample was subjected to hardness measurement with a Vickers hardness meter while changing the depth from the surface. The results are shown in Table 7 and FIG.

Table 7

  Referring to Table 7 and FIG. 8, in the sample WC-20% Co-0.7% Cr-0.4% V using WC powder having an average particle size of 0.55 μm belonging to the ultrafine particle system, The surface layer hardness reaches HV hardness 2050, and the effect of diffusion heat treatment is recognized.

  BN-coated WC-10% Co and WC-20% Co both have a diffusion depth of 3 to 4 mm, which is smaller than that of Example 1, and the surface layer hardness is also low. I understand. This is presumably because h-BN is a high-temperature stable compound and thus the reaction with the metal phase is difficult to proceed.

Here, the metal vapor deposition coating process, the boron trichloride is a metal chloride [BCl _3], methane [CH _4], examples will be described using hydrogen [H _2].
The CVD apparatus shown in FIG. 9 was used. Boron trichloride [BCl ₃ ], methane [CH ₄ ], and hydrogen [H ₂ ] gas cylinders 11, 12, and 13 are supplied to the heating furnace 1 through the flow meter 3 and the regulating valve 4. . A water ring pump 2 is connected to the heating furnace 1 so that the inside of the heating furnace can be set to a desired reduced pressure. Two types of sintered bodies used in Example 3 were set in the heating furnace 1, and CVD treatment was performed under the chemical vapor deposition conditions shown in the following table. When checking B _{4 C} deposition thickness of the sintered body surface after treatment, it was approximately 12～15Myuemu.
In this embodiment, the low-pressure CVD process is used, but in order to further increase the film thickness, a thermal CVD method or a laser CVD method may be used, and a desired coating layer thickness can be obtained.

上記被膜層は上記実施例３〜５と同様の熱処理により、所定の拡散熱処理効果が認められた。 Table 8 B ₄ C deposition conditions

The coating layer had a predetermined diffusion heat treatment effect by the same heat treatment as in Examples 3 to 5.

Cemented carbides used in general warm and hot regions have a WC average particle size of 3 μm or more, so we evaluated them using so-called medium to coarse WC powders.
WC-13% Co-2% Ni-1% Cr [15LB] and WC using a commercially available average particle size of 5.7 μm WC powder, 1.3 μm Co powder, 1.5 μm Ni powder, and Cr-C powder. Prepared and mixed in -18% Co-4% Ni-1.5% Cr [22HB] composition. A compacted body having the same shape as in Example 1 was produced from the obtained mixed powder, and then liquid phase sintering was performed at 1380 ° C. × 1 Hr in vacuum to obtain respective sintered materials.
Next, a coating material was prepared using silicon carbide SiC as a silicon source for heat treatment. The adjustment method was the same as in Example 1, and a 15% SiC-containing ethanol coating was prepared. The surface of the sintered material was coated by a dipping method, dried and subjected to diffusion heat treatment. The heat treatment temperature was 1300 ° C x 3 Hr. In addition, comparative evaluation was also performed on the sample as it was without the coating treatment.
The heat-treated sample was cut at a position of 15 mm in length, the cut surface was polished, the cross-sectional structure was observed, and then the hardness was measured with a Vickers hardness meter while changing the depth from the surface.
As a result of the structure observation, an improvement in the distribution density of WC particles was observed up to a depth of about 2 mm, and the structure of the structure was clearly rich in binder metal.
The results of the hardness measurement are shown in Table 9 and FIG.

図１０から明らかなように、粗粒 WC を使用しているため、硬さとしては比較的低い値であるが、内質部と比較すると、表層部硬さは顕著な増大が認められた。
又、ケイ素の拡散深さは、硬さ傾斜部と見なすと、ホウ素拡散素材よりも小さく、これはホウ素とケイ素の元素特性の違いによるものと考えられた。しかしながら、バインダ金属の拡散移動はホウ素と同様の挙動を示すことが確認され、温間・熱間工具に致命的なヒートクラックの抑制に対する表面圧縮残留応力の効果、並びに耐熱性、耐酸化性が付与されることは、高温領域に適用される工具として極めて有用な特徴を有するものである。
又、被覆材として、SiB₆ を使用すれば、ホウ素とケイ素の両特性が複合した表層部特性が得られる。 Table 9

As apparent from FIG. 10, since coarse WC is used, the hardness is a relatively low value, but the surface layer hardness is remarkably increased as compared with the inner part.
Further, the diffusion depth of silicon is considered to be smaller than that of the boron diffusion material when it is regarded as a hardness gradient portion, which is considered to be due to the difference in elemental characteristics between boron and silicon. However, it has been confirmed that the diffusion movement of the binder metal shows the same behavior as boron, and the effect of surface compressive residual stress on the suppression of heat cracks that are fatal to warm and hot tools, as well as heat resistance and oxidation resistance. Being imparted has extremely useful characteristics as a tool applied to a high temperature region.
Further, when SiB ₆ is used as a coating material, surface layer characteristics in which both characteristics of boron and silicon are combined can be obtained.

[performance test]
Sample preparation Commercially available WC powder and Co powder with an average particle size of 1.5 μm were weighed and blended into a WC-14% Co composition, inserted into a stainless steel pot together with ethanol solvent and cemented carbide balls, and pulverized and mixed for 30 hours. The obtained raw material slurry was put into a stirrer and the solvent was dried. Then, 1.5 wt% paraffin wax was added, and heated to 70 ° C. to prepare a finished powder. Similarly, commercially available WC powder and Co powder with an average particle size of 3.2 μm were weighed and blended into a WC-17% Co composition, and milled, dried and mixed with wax to produce a finished powder.
Next, using a φ25 mm press mold, the finished powder was filled into the mold cavity and pressed at a pressure of 1 ton / cm ² to produce a φ25 × 30 L mm compact. .
The obtained green compact was degreased and pre-sintered at 900 ° C. in a pre-sintering furnace and then subjected to a gradient treatment (PD). Some pre-sintered bodies were vacuum-sintered at 1350 ° C. to form sintered bodies and then subjected to a gradient treatment (SG). In addition, a sintered body of WC-17% Co alloy using WC powder of 3.2 μm was prepared and subjected to a gradient treatment (VG) under almost the same conditions.

Inclination treatment Here, # 200-B ₄ C powder was used as a diffusion material. Ethanol and B ₄ C powder are pulverized and mixed for 5 hours with a ball mill, and a B ₄ C coating material prepared by PEI is prepared. A predetermined amount is applied to the outer surface of the pre-sintered body and sintered body to be tilted. It dried and the inclination process was performed on various conditions shown in Table 10. The gradient-treated alloy thus obtained was cut and polished at the center of each sample, and characteristics such as structure observation, element concentration analysis, and hardness measurement were confirmed.

Table 10 WC (1.5μ) -14% Co Tilting condition

Tissue characteristics Samples PD125 and PD130 have clear “nests” that appear as dispersed black spots, and include internal defects as an alloy material. When an alloy tool is made of such a material, the “nest” becomes a starting point of destruction, so that it is apparent that the tool breaks in a very short time after the start of use.
Furthermore, in PD135 and PD140, which have increased the gradient treatment temperature, there is almost no “nest” that is an internal defect due to complete sintering densification, but the concentration gradient of the Co-bonded phase is very unclear from the surface to the interior. It has become a thing. This seems to be because the liquid phase appears in the entire base material, so that the concentration of the liquid phase is made uniform in the range from the surface B diffusion region to the internal non-diffusion region. There is no difference in the WC particle size between the surface layer and the inside.
On the other hand, in the SG120, SG125, and SG130 that have been subjected to the gradient treatment from the sintered body, no “nest” as an internal defect is recognized. As the gradient structure, the concentration gradient of the Co bonded phase from the surface layer to the inside can be confirmed very clearly. In this way, the gradient treatment is performed in contrast to the case where the gradient treatment is performed from the pre-sintered body and the case where the gradient treatment is performed from the sintered body, and the gradient treatment is performed below the liquid phase appearance temperature of the sintered base material. Is important. Note that no grain growth structure is observed even when the pre-sintered body is inclined.

Hardness characteristics Fig. 11 shows the hardness distribution by HV measurement from the surface layer to the inside. Since the measured values of PD125 and PD130 in which tissue defects were recognized varied, the description was omitted as data.
First, in the inclination treatment from the pre-sintered body, PD135 and PD140 show an improvement in surface hardness of about HV = 300 compared to the internal hardness of the base metal. This is considered to be a synergistic effect of improving hardness by reducing the amount of Co-bonded phase by about 3% in the surface layer, and improving hardness by solid solution strengthening and precipitation strengthening of the diffusion element B. SG125. Compared with the surface hardness by 130, the hardness is about HV = 200-300.
In the gradient treatment from the pre-sintered body, the B and Si elements used in the present invention, especially the B element, have a small activation energy and a high diffusion rate, so that diffusion proceeds rapidly in the presence of a liquid phase. For this reason, it does not become the state concentrated in the surface layer part, and does not contribute so much to remarkable solid solution strengthening and precipitation strengthening.
On the other hand, in SG120 to SG130 in which the gradient treatment is performed from the sintered body, a remarkable improvement in the surface hardness is recognized as a whole. As the tilting process temperature increases, the slope of the tilted area tends to increase. Incidentally, if the gradient treatment temperature is made higher than this, the surface hardness is lowered to the same level as PD140 because the liquid phase appears in the whole material in the 1400 ° C treatment, for example.

Co concentration comparison and HV-Co correlation Figure 12 shows the concentration distribution of Co by EDAX analysis from the surface layer to the inside. Inclination treatment from pre-sintered body: Co concentration distribution of PD135 and PD140 increases from the surface to the inside, but is very gentle, and the surface / internal concentration ratio bs / bi is PD135 = 0.66, PD140. = 0.87.
On the other hand, SG120, SG125, and SG130 according to the present invention have a remarkably small Co concentration on the surface, and show a sharp increasing tendency at a position near 2 mm from the surface. The bs / bi calculated in the same manner as described above is SG120 = 0.54, SG125 = 0.39, SG130 = 0.28, and is characterized by being extremely small.

Fracture toughness evaluation of surface layerFurthermore, in the present invention, because of the structure of a hard surface layer with a significantly reduced amount of binder phase and an interior with an increased amount of binder phase, the inclined surface layer has a large compression. Residual stress is generated. An example of these is shown from the fracture toughness evaluation by IF method.
This indicates a crack propagated from the HV indentation in the surface layer, but the crack length in the internal direction from the surface of the gradient structure was much shorter than the crack length in the vertical direction. This suggests that the inclined structure according to the present invention imparts an effective compressive residual stress to the surface layer, so that the fracture from the surface to the internal direction is unlikely to occur, which is contradictory to high hardness and high toughness. It shows that it has characteristics.

Summarizing the above results, B ₄ C was selected as the metalloid element from B, Si, P, and in particular, B ₄ C was applied, and various evaluations were performed. As a result, the following was found. .
1) In this invention, since an inclination process is performed from a sintered compact, an internal defect does not arise.
2) In the gradient treatment of the present invention, a hardness gradient of about HV = 400 to 500 is obtained.
3) In the gradient treatment of the present invention, a gradient structure can be obtained regardless of the WC grain size.
4) In the gradient treatment of the present invention, a gradient structure can be obtained by significantly reducing the binder phase concentration in the surface layer.
5) In the gradient treatment of the present invention, the growth of WC particles does not occur, and a gradient structure can be obtained regardless of the particle size control.
6) Since the compressive residual stress is generated in the surface layer in the gradient treatment of the present invention, the fracture toughness of the surface layer is greatly improved.

The front view which shows the helical gear in which a screw part has a gentle spiral shape. The front view which shows the metal mold | die of a helical gear. The front view which shows the excavation tool by which the cemented carbide was brazed to the S55C support metal fitting. A sintered body obtained by dipping in a 9% coating solution of B ₄ C, using a fine hard particle (particle size of 1 to 2 μm), covering, and heat-treating the sintered tool according to the embodiment of the present invention. (A) shows an internal part and (B) shows a surface layer part, respectively. Sintered body obtained by dipping in a 9% coating solution of B ₄ C using a coarse hard particle (particle size: 3 to 6 μm), covering, and heat-treating the sintered tool according to the embodiment of the present invention. (A) shows an internal part and (B) shows a surface layer part, respectively. The figure which shows the change of the hardness in the depth direction from the surface of the sintered compact manufactured by the manufacturing method which concerns on Example 3 of this invention. The figure which shows the change of the hardness in the depth direction from the surface of the sintered compact which concerns on another Example 4. FIG. Furthermore, the figure which shows the change of the hardness in the depth direction from the surface of the sintered compact concerning another Example 5. FIG. Schematic of the CVD apparatus which forms a film layer. The figure which shows the change of the hardness in the depth direction from the surface of the sintered compact manufactured by the manufacturing method which concerns on Example 6 of this invention. The graph which shows distribution of hardness by HV measurement from a surface layer part to an inside. The graph which shows distribution of Co density | concentration by EDAX analysis from a surface layer part to the inside. Photomicrograph showing results of fracture toughness evaluation test by IF method.

Claims (13)

M ₁₂ C type to M ₃ C type double carbide (M is one or more of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and one or more of Fe, Co, Ni) WC-Co compacted green body with the main component of the surface layer part being carburized, followed by liquid phase sintering and increasing the liquid phase sintering temperature using the liquid phase sintering temperature as an index, A method for producing a cemented carbide material, characterized in that the average particle size of the surface layer WC is increased. M ₁₂ C type double carbide (M is at least one of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and at least one of Fe, Co, Ni) The WC-Co compacted green body, which is the main component of the part, is treated at a carburizing temperature of 600 to 900 ° C. and a liquid phase sintering temperature of 1300 ° C. or higher. The manufacturing method of Claim 1 which forms the structure | tissue inclination smaller than it. Surface layer of M ₃ C type double carbide (M represents one or more of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and one or more of Fe, Co, Ni) The WC-CO green compact molded as the main component is treated at a carburizing temperature of 800 to 1100 ° C and a liquid phase sintering temperature of 1350 ° C or higher. The manufacturing method according to claim 1, wherein the structure is formed to have a larger tissue inclination.   The manufacturing method according to claim 1, wherein a compound containing boron or silicon is applied to the surface of the obtained sintered body, and diffusion heat treatment is performed in a temperature range of 1200 to 1350 ° C. below the liquidus temperature. M ₁₂ C type to M ₃ C type double carbide (M is one or more of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and one or more of Fe, Co, Ni) After the liquid phase sintering of the WC-Co green compact with the main component of the surface layer portion, a compound containing boron or silicon as melting point lowering elements is applied to the surface of the sintered body, and again A method for producing a sintered tool of high-strength cemented carbide, characterized by performing diffusion heat treatment in a temperature range of 1200 to 1350 ° C. below the liquid phase sintering temperature. The manufacturing method according to claim 5, wherein boron in the boron coating layer is in the range of 5.0 to 40 mg / cm ² with respect to the coating surface in terms of metallic boron element. M ₁₂ C type double carbide (M is at least one of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and at least one of Fe, Co, Ni) WC-Co-based sintered body as a main component of the surface part, and the surface layer part WC average grain size has a structure inclination that is 0.3 to 0.7 times smaller than that of the internal part, and the surface layer part A high-strength cemented carbide sintered tool having a concentration gradient in which the bonded metal is moved inward. The high-strength cemented carbide sintered tool according to claim 7, having excellent mechanical properties such that the surface layer hardness is HRA = 91 to 95 and the toughness is K _IC = 15 to 23 MN / m ^3/2 . Surface layer of M ₃ C type double carbide (M represents one or more of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and one or more of Fe, Co, Ni) WC-Co-based sintered body with the main component of the surface part, and the surface layer part WC average grain size is 1.5 times greater than that of the inner part part, and the surface layer part has a bonded metal on the inner side. A high-strength cemented carbide sintered tool with a shifted concentration gradient. The high-strength cemented carbide sintered tool according to claim 9, which has excellent mechanical properties such that the surface layer hardness is HRA = 88 to 92 and the toughness is K _IC = 20 to 30 MN / m ^3/2 . M ₁₂ C type to M ₃ C type double carbide (M is at least one of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, W and any of iron group metals Fe, Co, Ni) WC-Co-based sintered tool having a main component of a surface layer portion, the surface layer portion including boron B or silicon Si in a range of 0.010 to 1.0% by weight. The high-strength cemented carbide sintered tool, wherein the surface layer portion has hard particles having a distribution density higher than that of the internal portion.   The weight ratio of the content of the iron group metal (one or more of Fe, Co, Ni) and the hard particles WC in the inner part is in the range of 5:95 to 40:60. The high-strength cemented carbide sintered tool described. The sintered tool according to claim 11, wherein the binder metal content in the range from the surface of the surface layer portion to a depth of 0.5 mm is 2% or less by weight.

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