JP6924874B2 - Cobalt-based alloy material - Google Patents

Description

The present invention relates to a cobalt-based alloy material having excellent mechanical properties, and more particularly to an alloy material for obtaining the cobalt-based alloy material.

Cobalt (Co) -based alloy material is a typical heat-resistant alloy material together with nickel (Ni) -based alloy material, and is also called a superalloy, and is widely used for high-temperature members of turbines (for example, gas turbines and steam turbines). There is. Co-based alloy materials have been used as turbine stationary blades and combustor members because they have higher material cost than Ni-based alloy materials, but have excellent corrosion resistance and wear resistance, and are easily solid-solved and strengthened.

Due to various alloy composition improvements and manufacturing process improvements that have been made to date in heat-resistant alloy materials, Ni-based alloy materials are _{strengthened by precipitation of the γ'phase (for example, Ni 3} (Al, Ti) phase). It was developed and is now mainstream. On the other hand, in the Co-based alloy material, since it is difficult to precipitate an intermetallic compound phase such as the γ'phase of the Ni-based alloy material, which greatly contributes to the improvement of mechanical properties, precipitation strengthening by the carbide phase has been studied.

For example, in Patent Document 1 (Japanese Patent Laid-Open No. 61-243143), massive and granular carbides having a particle size of 0.5 to 10 μm are precipitated on a base of a cobalt-based alloy having a crystal particle size of 10 μm or less. A Co-based superplastic alloy characterized by the above is disclosed. The cobalt-based alloy has a weight ratio of C: 0.15 to 1%, Cr: 15 to 40%, W and / or Mo: 3 to 15%, B: 1% or less, Ni: 0 to 20%, Nb: It is disclosed that it consists of 0 to 1.0%, Zr: 0 to 1.0%, Ta: 0 to 1.0%, Ti: 0 to 3%, Al: 0 to 3%, and the balance Co. According to Patent Document 1, Co-based superplasticity that exhibits superplasticity even in a low temperature region (for example, 950 ° C.), has an elongation rate of 70% or more, and can produce a complicated shape by plastic working such as forging. It is said that it can provide alloys.

According to Patent Document 2 (Japanese Patent Laid-Open No. 7-179967), in% by weight, Cr: 21 to 29%, Mo: 15 to 24%, B: 0.5 to 2%, Si: 0.1% or more and less than 0.5%, C : A Co-based alloy having excellent corrosion resistance, wear resistance, and high-temperature strength, which is more than 1% and 2% or less, Fe: 2% or less, Ni: 2% or less, and the balance substantially Co, is disclosed. According to Patent Document 2, the Co-based alloy has a composite structure in which molybdenum boride and chromium carbide are relatively finely dispersed in a quaternary alloy phase of Co, Cr, Mo, and Si, and has good corrosion resistance and resistance. It is said to have abrasion resistance and high strength.

By the way, in recent years, three-dimensional modeling technology (so-called 3D printing) such as additive manufacturing (AM method) has attracted attention as a technology for manufacturing a final product having a complicated shape by a near net shape, and the three-dimensional modeling technology has been attracting attention. Is being actively researched and developed to apply the above to heat-resistant alloy members.

For example, Patent Document 3 (Japanese Patent Laid-Open No. 2016-535169) provides a layer forming method including the following steps: a) a raw material for a powdery or suspension-like granular composite material having a void ratio of less than 20%. Steps to b) deposit the first part of the composite on the surface of the target, c) supply energy to the composite in the first part to sinter, fuse or blend the composite in the first part. The step of melting to form the first layer, d) the step of depositing the second portion of the composite material on the first layer, e) supplying energy to the composite material of the second part, said Disclosed is a step of sintering, fusing or melting the composite material in a second portion to form a second layer, wherein the energy is supplied by a laser.

According to Patent Document 3 (Special Table 2016-535169), selective laser melting (SLM) or direct metal laser melting (DMLM) generally involves one type of material (pure titanium or Ti-6Al-4V). It is said to be beneficial for such single alloys). In contrast, for multiple materials, alloys, or combinations of alloys with other materials such as plastics, ceramics, polymers, carbides, and glass, selective laser sintering (SLS) or direct metal laser sintering (DMLS). The law) is generally applied. Sintering is a technical concept different from melting, and the sintering process is a process in which the material powder is not completely melted, but is heated to the point where the powder can be fused together at the molecular level. That is.

Japanese Unexamined Patent Publication No. 61-243143 Japanese Unexamined Patent Publication No. 7-179967 Special Table 2016-535169

Manufacture of turbine high-temperature members by 3D printing is very difficult from the viewpoint of shortening the manufacturing work time and improving the manufacturing yield (that is, from the viewpoint of reducing manufacturing costs) because even members with complicated shapes can be directly modeled. It's an attractive technology.

Co-based alloy materials as described in Patent Documents 1 and 2 are considered to have higher mechanical properties than the previous Co-based alloy materials, but are compared with recent precipitation-strengthened Ni-based alloy materials. Unfortunately, it cannot be said that it has sufficient mechanical properties. Therefore, at present, most of the research on laminated molded bodies (AM bodies) for turbine high-temperature members is targeted at precipitation-strengthened Ni-based alloy materials.

However, in the AM body of the precipitation-strengthened Ni-based alloy, problems such as inhibition of the formation of the γ'phase, which is the key to mechanical properties, and internal defects in the product are likely to occur, which is expected as a result. There is a problem that sufficient mechanical properties cannot be obtained. This is because the current precipitation-hardened Ni-based alloy material used as a high-temperature member for turbines is optimized on the premise of melting and casting processes in high vacuum, so it is necessary to prepare alloy powder for the AM method. It is considered that oxidation and nitriding of the Al component and Ti component constituting the γ'phase are likely to occur at the stage of the AM method.

On the other hand, the Co-based alloy material as described in Patent Documents 1 and 2 does not presuppose the precipitation of an intermetallic compound phase such as the γ'phase of the Ni-based alloy material, and therefore easily oxidizes Al and Ti. It does not contain much and can be used for melting and casting processes in the air. Therefore, it is considered to be advantageous for the production of alloy powders for the AM method and the production of AM bodies. Further, the Co-based alloy material has an advantage of having corrosion resistance and wear resistance equal to or higher than that of the Ni-based alloy material.

However, as described above, the conventional Co-based alloy material has a weakness that the mechanical properties are lower than those of the γ'phase precipitation strengthened Ni-based alloy material. In other words, achieve mechanical properties equal to or higher than those of the γ'phase precipitation strengthened Ni-based alloy material (for example, a creep withstand temperature of 875 ° C or higher for 100,000 hours at 58 MPa and a tensile strength at room temperature of 500 MPa or higher). If possible, the Co-based alloy AM body can be a very attractive turbine high temperature member.

The present invention has been made in view of the above problems, and an object of the present invention is to realize a Co-based alloy material having mechanical properties equal to or higher than those of a precipitation-strengthened Ni-based alloy material. An object of the present invention is to provide a Co-based alloy material suitable for producing a Co-based alloy laminated model having a structure.

(I) One aspect of the present invention is a Co-based alloy material.
With carbon (C) of 0.08% by mass or more and 0.25% by mass or less,
Boron (B) of 0.1% by mass or less and
Contains 10% by mass or more and 30% by mass or less of chromium (Cr)
It contains 5% by mass or less of iron (Fe) and 30% by mass or less of nickel (Ni), and the total of Fe and Ni is 30% by mass or less.
It contains tungsten (W) and / or molybdenum (Mo), and the total of the W and the Mo is 5% by mass or more and 12% by mass or less.
It contains titanium (Ti), zirconium (Zr), niobium (Nb) and tantalum (Ta), and the total of the Ti, the Zr, the Nb and the Ta is 0.5% by mass or more and 2% by mass or less.
With 0.5% by mass or less of silicon (Si),
With 0.5% by mass or less of manganese (Mn),
Containing with nitrogen (N) of 0.003% by mass or more and 0.04% by mass or less,
The balance has a chemical composition consisting of Co and impurities,
The impurities include 0.04% by mass or less of oxygen (O).
It provides a cobalt-based alloy material characterized by the above.

The present invention can make the following improvements and modifications to the above-mentioned Co-based alloy material (I).
(I) The chemical composition is
The Ti is 0.01% by mass or more and 1% by mass or less.
The Zr is 0.05% by mass or more and 1.5% by mass or less.
The Nb is 0.02% by mass or more and 1% by mass or less.
The Ta is 0.05% by mass or more and 1.5% by mass or less.
(Ii) The chemical composition further contains 0.5% by mass or less of aluminum (Al) as the impurities.
(Iii) The chemical composition further contains rhenium (Re) of 2% by mass or less.
The total of the W, the Mo and the Re is 5% by mass or more and 12% by mass or less.
(Iv) The Co-based alloy material is an alloy powder for a predetermined selective laser melting method.
(V) The particle size of the alloy powder is 5 μm or more and 100 μm or less.

According to the present invention, a Co group suitable for producing a Co-based alloy laminated model having a predetermined fine structure for realizing a Co-based alloy material having mechanical properties equal to or higher than that of a precipitation-strengthened Ni-based alloy material. Alloy materials can be provided.

It is a flow chart which shows the process example of the manufacturing method of the Co-based alloy product which concerns on this invention. It is a scanning electron microscope (SEM) observation image which shows an example of the microstructure of the Co-based alloy AM body which concerns on this invention. It is an example of the Co-based alloy product which concerns on this invention, and is the perspective schematic diagram which shows the turbine vane as a turbine high temperature member. It is sectional drawing which shows an example of the gas turbine equipped with the Co-based alloy product which concerns on this invention. It is an SEM observation image which shows an example of the microstructure of the Co-based alloy AM body formed by the laser metal deposition method (LMD method). It is an SEM observation image which shows an example of the microstructure of a Co-based alloy cast body formed by a precision casting method. It is a graph which shows the relation example of the average size of the segregation cell in a Co-based alloy AM body, and 0.2% proof stress in a Co-based alloy product. It is an example of SLM condition for obtaining a Co-based alloy AM body which concerns on this invention, and is a graph which shows the relationship between the thickness of an alloy powder bed, and the amount of local heat input.

[Basic Thought of the Present Invention]
As described above, in the Co-based alloy material, various research and development have been carried out for strengthening by precipitation of the carbide phase. Examples of the carbide phase that contributes to precipitation strengthening include MC-type carbide phases of Ti, Zr, Nb, and Ta, and composite carbide phases of these metal elements.

The C component, which is indispensable for forming the carbide phase with each component of Ti, Zr, Nb, and Ta, is remarkably located in the final solidified part (for example, the dendrite boundary and the grain boundary) during melt solidification of the Co-based alloy. Has the property of segregating. Therefore, in the conventional Co-based alloy material, the carbide phase particles are precipitated along the dendrite boundary of the matrix phase and the grain boundary. For example, in ordinary castings of Co-based alloys, the average spacing and average crystal grain size of dendrite boundaries are usually on the ^{order of 10 1} to 10 ² μm, so the average spacing of carbide phase particles is also on the order of ^{10 1} to 10 ^{2 μm.} Become. Further, even in a process such as laser welding in which the solidification rate is relatively high, the average spacing of the carbide phase particles in the solidification portion is about 5 μm.

It is generally known that precipitation strengthening in alloys is inversely proportional to the average spacing between precipitates, and it is said that precipitation strengthening is effective when the average spacing between precipitates is about 2 μm or less. There is. However, in the above-mentioned conventional technique, the average spacing between the precipitates does not reach that level, and a sufficient effect of precipitation strengthening cannot be obtained. In other words, in the prior art, it was difficult to finely disperse and precipitate carbide phase particles that contribute to alloy strengthening. This is the main reason why Co-based alloy materials have been said to have insufficient mechanical properties compared to precipitation-strengthened Ni-based alloy materials.

As another carbide phase that can be precipitated in the Co-based alloy, there is a Cr carbide phase. Since the Cr component has high solid solubility in the Co-based alloy matrix and is difficult to segregate, the Cr carbide phase can be dispersed and precipitated in the matrix crystal grains. However, it is known that the Cr carbide phase has low lattice consistency with the Co-based alloy matrix crystal and is not so effective as a precipitation strengthening phase.

The present inventors can dramatically improve the mechanical properties of the Co-based alloy material if the carbide phase particles that contribute to precipitation strengthening can be dispersed and precipitated in the parent phase crystal grains in the Co-based alloy material. I thought I could do it. In addition, it was considered that a heat-resistant alloy material superior to the precipitation-strengthened Ni-based alloy material could be provided in combination with the good corrosion resistance and wear resistance originally possessed by the Co-based alloy material.

Therefore, the present inventors have diligently studied the alloy composition and the manufacturing method for obtaining such a Co-based alloy material. As a result, in addition to optimizing the alloy composition, in the production using the AM method (particularly, the selective laser melting method), the Co group is controlled by controlling the amount of heat input for local melting and quenching solidification within a predetermined range. It has been found that carbide phase particles that contribute to alloy strengthening can be dispersed and precipitated in the matrix crystal grains of the alloy material. The present invention has been completed based on this finding.

Hereinafter, embodiments according to the present invention will be described with reference to the drawings along with manufacturing procedures. However, the present invention is not limited to the embodiments taken up here, and can be appropriately combined with a known technique or improved based on the known technique without departing from the technical idea of the invention. ..

[Manufacturing method of Co-based alloy laminated model / Co-based alloy product]
FIG. 1 is a flow chart showing a process example of a method for producing a Co-based alloy product according to the present invention. As shown in FIG. 1, the method for producing a Co-based alloy product according to the present invention generally uses an alloy powder preparation step (S1) for preparing a Co-based alloy powder and the prepared Co-based alloy powder. A selective laser melting step (S2) for forming an AM alloy having a desired shape, a solution heat treatment step (S3) for subjecting the formed AM alloy to a solution treatment, and an AM alloy having a solution treatment. It has an aging heat treatment step (S4) of performing an aging treatment. The AM body obtained by the selective laser melting step S2 is the Co-based alloy laminated model according to the present invention.

Hereinafter, each step will be described in more detail.

(Alloy powder preparation process)
This step S1 is a step of preparing a Co-based alloy powder having a predetermined chemical composition. The chemical composition consists of C of 0.08% by mass or more and 0.25% by mass or less, B of 0.1% by mass or less, Cr of 10% by mass or more and 30% by mass or less, and Fe of 5% by mass or less and a total of 30% by mass or less. Fe and Ni, W and / or Mo with a total of 5% by mass or more and 12% by mass or less, Ti, Zr, Nb and Ta with a total of 0.5% by mass or more and 2% by mass or less, and Si with a total of 0.5% by mass or less. It is preferable that Mn of 0.5% by mass or less and N of 0.003% by mass or more and 0.04% by mass or less are contained, and the balance is composed of Co and impurities.

C: 0.08% by mass or more and 0.25% by mass or less
The C component is an important component constituting the MC-type carbide phase (which may be referred to as a carbide phase of Ti, Zr, Nb and / or Ta, or a strengthened carbide phase) to be a precipitation strengthening phase. The content of the C component is preferably 0.08% by mass or more and 0.25% by mass or less, more preferably 0.1% by mass or more and 0.2% by mass or less, and further preferably 0.12% by mass or more and 0.18% by mass or less. If the C content is less than 0.08% by mass, the amount of the reinforced carbide phase precipitated is insufficient, and the effect of improving the mechanical properties cannot be sufficiently obtained. On the other hand, when the C content exceeds 0.25% by mass, the alloy material is excessively hardened, and the ductility and toughness of the alloy material are lowered.

B: 0.1% by mass or less
The B component is a component that contributes to the improvement of the bondability of the crystal grain boundaries (so-called grain boundary strengthening). The component B is not an essential component, but when it is contained, it is preferably 0.1% by mass or less, more preferably 0.005% by mass or more and 0.05% by mass or less. When the B content exceeds 0.1% by mass, cracks (for example, solidification cracks) are likely to occur during the formation of the AM body.

Cr: 10% by mass or more and 30% by mass or less
The Cr component is a component that contributes to the improvement of corrosion resistance and oxidation resistance. The content of the Cr component is preferably 10% by mass or more and 30% by mass or less, and more preferably 10% by mass or more and 25% by mass or less. When a corrosion-resistant coating layer is separately provided on the outermost surface of the Co-based alloy product, the content of the Cr component is more preferably 10% by mass or more and 18% by mass or less. If the Cr content is less than 10% by mass, the corrosion resistance and oxidation resistance become insufficient. On the other hand, when the Cr content exceeds 30% by mass, a brittle σ phase is formed or a Cr carbide phase is formed, and the mechanical properties (toughness, ductility, strength) are lowered.

Ni: 30% by mass or less
Since the Ni component has properties similar to those of the Co component and is cheaper than the Co component, it is a component that can be contained in a form that replaces a part of the Co component. The Ni component is not an essential component, but when it is contained, it is preferably 30% by mass or less, more preferably 20% by mass or less, and further preferably 5% by mass or more and 15% by mass or less. When the Ni content exceeds 30% by mass, the wear resistance and resistance to local stress, which are the characteristics of Co-based alloys, decrease. This is considered to be due to the difference between the stacking defect energy of Co and that of Ni.

Fe: 5% by mass or less
Since the Fe component is much cheaper than the Ni component and has properties similar to those of the Ni component, it is a component that can be contained in a form that replaces a part of the Ni component. That is, the total content of Fe and Ni is preferably 30% by mass or less, more preferably 20% by mass or less, and further preferably 5% by mass or more and 15% by mass or less. The Fe component is not an essential component, but when it is contained, it is preferably 5% by mass or less, more preferably 3% by mass or less in a range smaller than the Ni content. When the Fe content exceeds 5% by mass, it becomes a factor of deterioration of corrosion resistance and mechanical properties.

W and / or Mo: Total 5% by mass or more and 12% by mass or less
The W component and the Mo component are components that contribute to the solid solution strengthening of the matrix. The total content of the W component and / or the Mo component is preferably 5% by mass or more and 12% by mass or less, and more preferably 7% by mass or more and 10% by mass or less. When the total content of the W component and the Mo component is less than 5% by mass, the solid solution strengthening of the matrix phase becomes insufficient. On the other hand, when the total content of the W component and the Mo component exceeds 12% by mass, the brittle σ phase is likely to be generated and the mechanical properties (toughness, ductility) are lowered.

Re: 2% by mass or less
The Re component is a component that contributes to the strengthening of the solid solution of the matrix and the improvement of corrosion resistance. The Re component is not an essential component, but when it is contained, it is preferably 2% by mass or less in the form of replacing a part of the W component or the Mo component, and more preferably 0.5% by mass or more and 1.5% by mass or less. When the Re content exceeds 2% by mass, the action and effect of the Re component is saturated, and the increase in material cost becomes a demerit.

Ti, Zr, Nb and Ta: Total 0.5% by mass or more and 2% by mass or less
The Ti component, Zr component, Nb component and Ta component are important components constituting the reinforced carbide phase (MC type carbide phase). The total content of the Ti, Zr, Nb and Ta components is preferably 0.5% by mass or more and 2% by mass or less, and more preferably 0.5% by mass or more and 1.8% by mass or less in total. If the total content is less than 0.5% by mass, the amount of the reinforced carbide phase precipitated is insufficient, and the effect of improving the mechanical properties cannot be sufficiently obtained. On the other hand, when the total content exceeds 2% by mass, the strengthened carbide phase particles are coarsened, the formation of brittle phase (for example, σ phase) is promoted, and oxide phase particles that do not contribute to precipitation strengthening are generated. Mechanical properties are reduced.

More specifically, the Ti content is preferably 0.01% by mass or more and 1% by mass or less, and more preferably 0.05% by mass or more and 0.8% by mass or less. The Zr content is preferably 0.05% by mass or more and 1.5% by mass or less, and more preferably 0.1% by mass or more and 1.2% by mass or less. The Nb content is preferably 0.02% by mass or more and 1% by mass or less, and more preferably 0.05% by mass or more and 0.8% by mass or less. The Ta content is preferably 0.05% by mass or more and 1.5% by mass or less, and more preferably 0.1% by mass or more and 1.2% by mass or less.

Si: 0.5% by mass or less
The Si component is a component that plays a role of deoxidizing and contributes to the improvement of mechanical properties. The Si component is not an essential component, but when it is contained, it is preferably 0.5% by mass or less, and more preferably 0.01% by mass or more and 0.3% by mass or less. When the Si content exceeds 0.5% by mass _{, coarse particles of oxide (for example, SiO 2} ) are formed, which causes deterioration of mechanical properties.

Mn: 0.5% by mass or less
The Mn component is a component that plays a role of deoxidation and desulfurization and contributes to improvement of mechanical properties and corrosion resistance. The Mn component is not an essential component, but when it is contained, it is preferably 0.5% by mass or less, more preferably 0.01% by mass or more and 0.3% by mass or less. When the Mn content exceeds 0.5% by mass, coarse particles of sulfide (for example, MnS) are formed, which causes deterioration of mechanical properties and corrosion resistance.

N: 0.003% by mass or more and 0.04% by mass or less
The N component is a component that contributes to the stable formation of the strengthened carbide phase. The content of the N component is preferably 0.003% by mass or more and 0.04% by mass or less, more preferably 0.005% by mass or more and 0.03% by mass or less, and further preferably 0.007% by mass or more and 0.025% by mass or less. If the N content is less than 0.003% by mass, the action and effect of the N component cannot be sufficiently obtained. On the other hand, when the N content exceeds 0.04% by mass, coarse particles of nitride (for example, Cr nitride) are formed, which causes a decrease in mechanical properties.

Remaining: Co component + impurities
The Co component is one of the main components of this alloy and is the component having the maximum content. As described above, the Co-based alloy material has an advantage of having corrosion resistance and wear resistance equal to or higher than that of the Ni-based alloy material.

The Al component is one of the impurities of this alloy and is not a component intentionally contained. However, if the Al content is 0.3% by mass or less, it is acceptable because it does not significantly adversely affect the mechanical properties of the Co-based alloy product. When the Al content exceeds 0.3% by mass _{, coarse particles of oxides and nitrides (for example, Al 2} O ₃ and Al N) are formed, which causes deterioration of mechanical properties.

The O component is also one of the impurities of this alloy and is not a component intentionally contained. However, an O content of 0.04% by mass or less is acceptable because it does not significantly adversely affect the mechanical properties of the Co-based alloy product. When the O content exceeds 0.04% by mass, coarse particles of various oxides (for example, Ti oxide, Zr oxide, Al oxide, Fe oxide, Si oxide) are formed, which causes deterioration of mechanical properties. become.

In this step S1, there is no particular limitation on the method / method for preparing the Co-based alloy powder, and the conventional method / method can be used. For example, a mother alloy ingot production element step (S1a) for producing a mother alloy ingot (master ingot) by mixing, melting, and casting raw materials so as to have a desired chemical composition, and forming an alloy powder from the mother alloy ingot. The atomizing elementary process (S1b) may be performed. In addition, there are no particular restrictions on the atomizing method, and conventional methods and methods can be used. For example, a gas atomizing method or a centrifugal atomizing method that can obtain high-purity, homogeneous composition, and spherical particles can be preferably used.

The particle size of the alloy powder is preferably 5 μm or more and 100 μm or less, more preferably 10 μm or more and 70 μm or less, and 10 μm or more and 50 μm or less from the viewpoint of handleability and filling property of the alloy powder bed in the selective laser melting step S2 of the next step. More preferred. If the particle size of the alloy powder is less than 5 μm, the fluidity of the alloy powder decreases in the next step S2 (the formability of the alloy powder bed decreases), which causes the shape accuracy of the AM body to decrease. On the other hand, when the particle size of the alloy powder exceeds 100 μm, it becomes difficult to control the local melting and quenching solidification of the alloy powder bed in the next step S2, the melting of the alloy powder becomes insufficient, and the surface roughness of the AM body increases. It becomes a factor to do.

From the above, it is preferable to carry out the alloy powder classifying step (S1c) for classifying the particle size of the alloy powder into the range of 5 μm or more and 100 μm or less. In the present invention, even if it is confirmed that the particle size distribution of the alloy powder produced in the atomizing elementary step S1b is within a desired range as a result of measuring the particle size distribution, it is considered that the present elementary step S1c has been performed.

(Selective laser melting process)
This step S2 is a step of forming an AM body having a desired shape by a selective laser melting (SLM) method using the prepared Co-based alloy powder. Specifically, an alloy powder bed preparation step (S2a) in which a Co-based alloy powder is spread to prepare an alloy powder bed having a predetermined thickness, and a predetermined region of the alloy powder bed is irradiated with laser light to prepare the alloy powder bed in the region. This is a process of forming an AM body by repeating the laser melt coagulant step (S2b) in which the Co-based alloy powder is locally melted and rapidly cooled and solidified.

In this step S2, in order to obtain a desired microstructure (microstructure in which reinforced carbide phase particles are dispersed and precipitated in the parent phase crystal grains) in the final Co-based alloy product, AM which is a precursor of the product is used. Controls the microstructure of the body. Then, in order to control the microstructure of the AM body, local melting / quenching solidification of the alloy powder bed is controlled.

More specifically, in relation to the thickness h (unit: μm) of the alloy powder bed, the output P (unit: W) of the laser beam, and the scanning speed S (unit: mm / s) of the laser beam, “15 The thickness h of the alloy powder bed, the laser light output P, and the laser light scanning speed S so as to satisfy <h <150 "and" 67 (P / S) -3.5 <h <2222 (P / S) +13 ". It is preferable to control. If the control conditions are not met, an AM body having a desired microstructure cannot be obtained. By this step S2, the Co-based alloy AM body according to the present invention is obtained.

The output P of the laser beam and the scanning speed S of the laser beam basically depend on the configuration of the laser device, but should be selected within the range of "10 ≤ P ≤ 1000" and "10 ≤ S ≤ 7000", for example. Just do it.

(Co-based alloy laminated model)
FIG. 2 is a scanning electron microscope (SEM) observation image showing an example of the microstructure of the Co-based alloy AM body according to the present invention. As shown in FIG. 2, the Co-based alloy AM body of the present invention has an extremely specific microstructure that has never been seen before.

The AM body is a polycrystal having an average crystal grain size of 10 μm or more and 100 μm or less, and segregated cells having an average size of 0.15 μm or more and 1.5 μm or less are formed in the crystal grains of the polycrystal. In addition, particles of the reinforced carbide phase are precipitated at an average interval of 0.15 μm or more and 1.5 μm or less. In the present invention, the size of the segregated cell is defined as the average of the major axis and the minor axis.

Further detailed microstructure observation was performed using a transmission electron microscope-energy dispersive X-ray spectroscopy (TEM-EDX). It was confirmed that the components (Ti, Zr, Nb, Ta, C) forming the strengthened carbide phase were segregated in the cell wall-like region. In addition, it was confirmed that particles of the reinforced carbide phase were precipitated at a part of the triple point / four emphasis in the boundary region of the segregation cell.

(Solution heat treatment process)
This step S3 is a step of subjecting the formed Co-based alloy AM body to a solution treatment. As the solution treatment conditions, heat treatment with a holding time of 0.5 hours or more and 10 hours or less in a temperature range of 1100 ° C. or higher and 1200 ° C. or lower is preferable. The cooling method after the heat treatment is not particularly limited, and any of water cooling, oil cooling, air cooling, and furnace cooling may be used.

By performing the solution treatment, recrystallization of the matrix crystal grains occurs in the AM body, and the internal strain of the AM body generated at the time of quenching solidification can be relaxed. Further, it is preferable to control the coarseness of the average crystal grain size of the matrix crystal grains in the range of 20 μm or more and 145 μm or less by recrystallization. If the average crystal grain size is less than 20 μm or more than 145 μm, sufficient creep characteristics cannot be obtained in the final Co-based alloy product.

In addition, very interestingly, during the recrystallization of the parent phase grains, the components segregated in the boundary region of the segregated cell earlier began to aggregate to form a strengthened carbide phase, resulting in the segregated cell. It was found to disappear (more accurately, the segregated cells could not be confirmed by SEM observation). The agglomeration points at which the strengthened carbide phase begins to form are considered to be the triple points / four-focused positions of the boundary of the original segregation cell, and are finely distributed throughout the parent phase crystal grains (inside the crystal grains and on the grain boundaries). Become.

If the temperature conditions and holding time conditions of the solution heat treatment are well controlled, the strengthened carbide phase that has begun to be formed can be grown to the form of particles without coagulation and coarsening. In that case, the Co-based alloy product may be completed by this step S3. However, from the viewpoint of controlling the coarsening of the matrix crystal grains (in other words, from the viewpoint of production stability and production yield), it is more preferable to perform the following aging heat treatment step S4.

(Aging heat treatment process)
This step S4 is a step of aging the Co-based alloy AM body that has been solution-treated. As the aging treatment conditions, heat treatment with a holding time of 0.5 hours or more and 10 hours or less in a temperature range of 750 ° C. or more and 1000 ° C. or less is preferable. The cooling method after the heat treatment is not particularly limited, and any of water cooling, oil cooling, air cooling, and furnace cooling may be used.

By applying the aging treatment, the strengthened carbide phase that has begun to be formed in the solution heat treatment step S3 can be grown to the particle form while suppressing the excessive coarsening of the parent phase crystal grains. As a result, the average crystal grain size of the parent phase crystal grains is 20 μm or more and 145 μm or less, the reinforced carbide phase particles are finely dispersed and precipitated in each crystal grain, and the average interparticle distance of the reinforced carbide phase particles is 0.15 μm. A Co-based alloy product having a size of 1.5 μm or more can be obtained. As a matter of course, in the Co-based alloy product of the present invention, the reinforced carbide phase particles are dispersed and precipitated on the grain boundaries of the parent phase crystal grains.

(Finishing process)
Although not shown in FIG. 1, a step of forming a corrosion-resistant coating layer and a surface finishing step of the Co-based alloy product obtained by the solution heat treatment step S3 and the aging heat treatment step S4, if necessary. May be further performed.

[Co-based alloy product]
FIG. 3 is an example of a Co-based alloy product according to the present invention, and is a schematic perspective view showing a turbine vane as a turbine high temperature member. As shown in FIG. 3, the turbine stationary blade 100 is roughly composed of an inner ring side end wall 101, a blade portion 102, and an outer ring side end wall 103. A cooling structure is often formed inside the wing. As described above, since the turbine vane 100 has a very complicated shape and structure, the technical significance of the AM body formed by the near net shape and the alloy product based on the AM body is great.

For example, in the case of a gas turbine for power generation with an output of 30 MW class, the length of the blade of the turbine stationary blade (distance between both end walls) is about 170 mm.

FIG. 4 is a schematic cross-sectional view showing an example of a gas turbine equipped with the Co-based alloy product according to the present invention. As shown in FIG. 4, the gas turbine 200 is roughly composed of a compressor unit 210 that compresses intake air and a turbine unit 220 that blows fuel combustion gas onto turbine blades to obtain rotational power. The turbine high temperature member of the present invention can be suitably used as a turbine nozzle 221 or a turbine stationary blade 100 in the turbine section 220. As a matter of course, the turbine high temperature member of the present invention is not limited to gas turbine applications, and may be used for other turbine applications (for example, steam turbine applications).

Hereinafter, the present invention will be described in more detail with reference to Examples and Comparative Examples. The present invention is not limited to these examples.

[Experiment 1]
(Preparation of alloy powders IA-1 to IA-5 and CA-1 to CA-5)
A Co-based alloy powder having the chemical composition shown in Table 1 was prepared (alloy powder preparation step S1). More specifically, first, the raw materials were mixed, and then melted and cast by a vacuum high-frequency induction melting method to prepare a mother alloy ingot (mass: about 2 kg). Next, the atomizing element step S1b was performed in which the mother alloy ingot was redissolved to form an alloy powder by a gas atomizing method in an argon gas atmosphere.

Next, the alloy powder classifying step S1c for controlling the particle size was performed on each of the obtained alloy powders. At this time, an alloy powder having a powder particle size in the range of 10 to 25 μm (particle size S) and an alloy powder having a powder particle size in the range of 100 to 150 μm (particle size L) were prepared.

As shown in Table 1, the alloy powders IA-1 to IA-5 are alloy powders having a chemical composition satisfying the provisions of the present invention. On the other hand, the alloy powder CA-1 has a C content and a Cr content outside the provisions of the present invention. In the alloy powder CA-2, the C content, the Ni content, and the total content of "Ti + Zr + Nb + Ta" are out of the scope of the present invention. In the alloy powder CA-3, the N content and the total content of "Ti + Zr + Nb + Ta" are out of the provisions of the present invention. The total content of "Ti + Zr + Nb + Ta" of the alloy powder CA-4 is out of the specification of the present invention. In the alloy powder CA-5, the W content and the total content of "Ti + Zr + Nb + Ta" are out of the provisions of the present invention.

[Experiment 2]
(Preparation of SLM alloy products using IA-2 powder and SLM alloy products using CA-5 powder)
Using the alloy powder of IA-2 and CA-5 particle size S prepared in Experiment 1, an AM body (diameter 8 mm × height 10 mm) was formed by the SLM method (selective laser melting step S2). The SLM conditions are that the thickness h of the alloy powder bed is 100 μm, the output P of the laser beam is 100 W, and the scanning speed S (mm / s) of the laser beam is variously changed to obtain the local heat input P / S (unit). : W · s / mm = J / mm) was controlled. Controlling the amount of local heat input corresponds to controlling the cooling rate.

Each AM body prepared above was heat-treated by holding it at 1150 ° C. for 4 hours (solution heat treatment step S3). Next, each AM body that had been solution-treated was heat-treated at 900 ° C. for 4 hours (aging heat treatment step S4), and the SLM alloy product using IA-2 powder and CA-5 powder were used. An SLM alloy product was prepared using the above.

(Preparation of LMD alloy products using IA-2 powder and LMD alloy products using CA-5 powder)
AM bodies are formed by the laser metal deposition method (LMD method) using alloy powders of IA-2 and CA-5 particle size L prepared in Experiment 1, and solution heat treatment steps S3 and aging heat treatment steps S4 similar to the above. To prepare an LMD alloy product using IA-2 powder and an LMD alloy product using CA-5 powder. The LMD conditions were that the output P of the laser beam was 800 W and the scanning speed S of the laser beam was 15 mm / s.

The LMD method is a kind of AM method, and is a method of laminating and forming by simultaneously performing alloy powder injection and laser irradiation. The amount of local heat input by the LMD method is generally larger than the amount of local heat input by the SLM method. In other words, the cooling rate in the LMD method is slower than the cooling rate in the SLM method.

(Manufacturing of cast alloy products using IA-2 powder and cast alloy products using CA-5 powder)
A cast body (diameter 8 mm x height 10 mm) was formed by a precision casting method using the alloy powder of IA-2 and CA-5 particle size L prepared in Experiment 1, and the same solution heat treatment step S3 as described above. And the aging heat treatment step S4 were carried out to prepare a cast alloy product using IA-2 powder and a cast alloy product using CA-5 powder.

(Microstructure observation and mechanical property test)
From the AM body, the cast body, and the product prepared above, test pieces for microstructure observation and mechanical property test were collected, respectively, and microstructure observation and mechanical property test were performed.

Microstructure observation was performed by SEM. In addition, by image analysis of the obtained SEM observation image using image processing software (ImageJ, public domain software developed by the National Institutes of Health (NIH)), the average size of segregated cells, the average interval of microsegregation, and the average interval of microsegregation, and The average interparticle distance of the carbide phase particles was measured.

As a mechanical property test, a tensile test was conducted in a room temperature environment (about 23 ° C.), and 0.2% proof stress was measured.

FIG. 5 is an SEM observation image showing an example of the fine structure of the Co-based alloy AM body formed by the LMD method, and FIG. 6 is an SEM showing an example of the fine structure of the Co-based alloy cast body formed by the precision casting method. It is an observation image. Further, FIG. 2 shown above is an SEM observation image showing an example of the fine structure of the Co-based alloy AM body formed by the SLM method. The samples in FIGS. 2 to 6 all use IA-2 as the alloy powder.

As described above, in the AM body (see FIG. 2) formed by the SLM method, segregated cells having a size of about 1 μm are formed in the crystal grains of the polycrystalline body. On the other hand, in the AM body (see FIG. 5) formed by the LMD method, each crystal grain of the polycrystalline body is composed of segregated cells having a size of about 5 to 20 μm. In the cast body formed by the precision casting method (see FIG. 6), microsegregation along the dendrite boundary is observed, and the interval between the microsegregations is about 100 to 300 μm.

When the microstructure of the product produced by dissolving and aging these AM bodies and cast bodies was observed and the average interparticle distance of the carbide phase particles was measured, the average interparticle distance was determined by the segregation cell. It was found that the average size and the average interval of microsegregation were almost the same (the illustration of the microstructure is omitted). In addition, when the average size of the segregation cell is very small (for example, when the average size is about 0.1 μm or less), the nearby carbide phases are coalesced by solution heat treatment and aging treatment to form large particles (as a result). The average interparticle distance of the carbide phase particles increases).

Since products with significantly different microstructures were obtained, the relationship between the average size of segregated cells and mechanical properties was investigated. FIG. 7 is a graph showing an example of the relationship between the average size of segregated cells in the Co-based alloy AM body and the 0.2% proof stress in the Co-based alloy product. Note that FIG. 7 also shows data on the cast body and the cast alloy product for comparison. In the cast, the average size of the segregation cells was substituted by the average interval of microsegregation.

As shown in FIG. 7, the Co-based alloy product prepared using the CA-5 powder showed an almost constant 0.2% proof stress regardless of the average size of the segregated cells. On the other hand, the Co-based alloy product prepared using the IA-2 powder had a large change in 0.2% proof stress depending on the average size of the segregated cells.

The total content of "Ti + Zr + Nb + Ta" in CA-5 powder is too low (almost not contained). Therefore, in the product using CA-5 powder, the reinforced carbide phase had a fine structure in which Cr carbide particles were precipitated without precipitation. From this result, it is confirmed that the Cr carbide particles are not so effective as the precipitation strengthening particles. On the other hand, the product using the IA-2 powder had a fine structure in which the strengthened carbide phase particles were precipitated. Therefore, it is considered that the 0.2% proof stress changed significantly depending on the average size of the segregated cells (the resulting average interparticle distance of the carbide phase particles).

Further, considering the required characteristics for the turbine high temperature member targeted by the present invention, a 0.2% proof stress of 500 MPa or more is required. Therefore, if 0.2% proof stress of 500 MPa or more is judged as "pass" and less than 500 MPa is judged as "fail", the average size of the segregated cells (the resulting average interparticle distance of the carbide phase particles) is 0.15. It was confirmed that mechanical properties that were "passed" were obtained in the range of ~ 1.5 μm. In other words, one of the reasons why sufficient mechanical properties could not be obtained in the conventional carbide phase-precipitated Co-based alloy material is considered to be that the average interparticle distance of the reinforced carbide phase particles could not be controlled within a desirable range.

[Experiment 3]
(SLM alloy products using IA-1 to IA-5 powder IP-1-1 to IP-5-1 and SLM alloy products using CA-1 to CA-5 powder CP-1-1 to Preparation of CP-5-1)
AM bodies (diameter 8 mm x height 10 mm) were formed by the SLM method using alloy powders of IA-1 to IA-5 and CA-1 to CA-5 particle size S prepared in Experiment 1 (selective). Laser melting process S2). As the SLM condition, based on the result of Experiment 2, the average interparticle distance of the carbide phase particles was controlled to be in the range of 0.15 to 1.5 μm.

Each AM body prepared above was heat-treated at 1150 ° C. for 4 hours (solution heat treatment step S3). Next, each AM alloy that has been solution-treated is heat-treated at 750 to 1000 ° C. for 0.5 to 10 hours (aging heat treatment step S4), and SLM using IA-1 to IA-5 powders is used. SLM alloy products CP-1-1 to CP-5-1 were prepared using alloy products IP-1-1 to IP-5-1 and CA-1 to CA-5 powders.

(Microstructure observation and mechanical property test)
From the SLM alloy products IP-1-1 to IP-5-1 and CP-1-1 to CP-5-1 prepared above, test pieces for microstructure observation and mechanical property test were collected, respectively. , Microstructure observation and mechanical property test were performed.

For microstructure observation, the same image analysis as in Experiment 2 was performed on the SEM observation image, and the average particle size of the matrix crystal grains and the average interparticle distance of the carbide phase particles were measured.

As a mechanical property test, a creep test was performed under the conditions of a temperature of 900 ° C. and a stress of 98 MPa, and the creep rupture time was measured. From the required characteristics for the turbine high temperature member targeted by the present invention, a creep rupture time of 1100 hours or more was determined as "pass", and a creep rupture time of less than 1100 hours was determined as "fail". A passing creep property means that the temperature at which the creep rupture time is 100,000 hours at a stress of 58 MPa is 875 ° C or higher. This creep characteristic can be said to be the same creep characteristic as the Ni-based alloy material.

The results of Experiment 3 are shown in Table 2.

As shown in Table 2, all the samples of the SLM alloy products IP-1-1 to IP-5-1 passed the creep test. This is because the average particle size of the matrix crystal grains is in an appropriate range, and the average particle distance of the strengthened carbide phase particles (MC type carbide phase particles of Ti, Zr, Nb and / or Ta) is sufficient. It is considered that this is due to the small size (that is, the strengthened carbide phase particles are finely dispersed and precipitated).

On the other hand, in the SLM alloy products CP-1-1 to CP-5-1, the average particle size of the parent phase crystal grains was in an appropriate range, but the creep test failed in all the samples. When viewed individually, CP-1-1 is considered to be due to the predominant precipitation of Cr carbide particles because the C content and Cr content are excessive. Since the C content and the total content of "Ti + Zr + Nb + Ta" are excessive in CP-2-1, it is considered that the average particle distance is increased as a result of the coarsening of the strengthened carbide phase particles. Since the C content of CP-3-1 is excessive and the total content of "Ti + Zr + Nb + Ta" is too small, it is considered that the Cr carbide particles were predominantly precipitated. From the results of CP-1-1 and CP-3-1, it is confirmed that Cr carbide particles are not so effective as precipitation strengthening particles. CP-4-1 and CP-5-1 are considered to be due to the fact that the strengthened carbide phase itself did not precipitate because the total content of "Ti + Zr + Nb + Ta" was too low (because it was hardly contained). Be done.

From the results of Experiment 3, it was confirmed that IA-1 to IA-5 having the chemical composition specified in the present invention are preferable as the starting powder of the SLM alloy product. It was also confirmed that the creep characteristics can be improved by controlling the average interparticle distance of the strengthened carbide phase particles in the range of 0.15 to 1.5 μm.

[Experiment 4]
(Preparation of SLM alloy products IP-1-2 to IP-1-7 and IP-2-2 to IP-2-7)
Using the alloy powder of IA-1 and IA-2 particle size S prepared in Experiment 1, an AM body (diameter 8 mm × height 10 mm) was prepared by the SLM method (selective laser melting step S2). As the SLM condition, based on the result of Experiment 2, the average interparticle distance of the carbide phase particles was controlled to be in the range of 0.15 to 1.5 μm.

Each AM body prepared above was subjected to solution heat treatment and aging treatment. At this time, by variously changing the temperature range (1000 to 1300 ° C.) and the holding time range (0.5 to 10 hours) of the solution treatment, the average particle size of the parent phase crystal grains is different, and the SLM alloy product IP-1- 2-IP-1-7 and IP-2-2 to IP-2-7 were prepared. The aging treatment conditions were the same as in Experiment 3.

(Microstructure observation and mechanical property test)
From the SLM alloy products IP-1-2 to IP-1-7 and IP-2-2 to IP-2-7 prepared above, test pieces for microstructure observation and mechanical property test were collected, respectively. , Microstructure observation and mechanical property test were performed.

For microstructure observation, image analysis was performed on the SEM observation image similar to Experiment 2, and the average particle size of the matrix crystal grains was measured. As the mechanical property test, the same creep test as in Experiment 2 was performed, and "pass / fail" was determined based on the same criteria as in Experiment 2. The results of Experiment 4 are shown in Table 3.

As shown in Table 3, it was confirmed that the average particle size of the matrix crystal grains is preferably 20 to 145 μm. Further, from the results of Experiment 4, it was confirmed that the solution treatment is preferably a heat treatment in which the solution treatment is maintained in a temperature range of 1100 to 1200 ° C. for 0.5 to 10 hours.

[Experiment 5]
(Examination of SLM conditions in selective laser melting process)
An AM body (diameter 8 mm × height 10 mm) was formed by the SLM method using the IA-4 particle size S alloy powder prepared in Experiment 1 (selective laser melting step S2). The SLM condition is that the output P of the laser beam is 85 W, and the local heat input P / S (unit: W. s / mm = J / mm) was controlled.

The average size of the segregated cells was measured by observing the microstructure of each AM body prepared above. As the microstructure observation / measurement method, SEM and image processing software (ImageJ) were used as in Experiment 2.

FIG. 8 is an example of SLM conditions for obtaining a Co-based alloy AM body according to the present invention, and is a graph showing the relationship between the thickness of the alloy powder bed and the amount of local heat input. In FIG. 8, as a result of observing the microstructure of the AM body, those having an average size of segregated cells in the range of 0.15 to 1.5 μm are judged as “passed” and indicated by “◯” in the figure, and other cells are indicated by “○”. It was judged as "failed" and indicated by "x" in the figure.

From the results of Experiment 5, the SLM conditions in the selective laser melting step S2 are the thickness h (unit: μm) of the alloy powder bed, the laser light output P (unit: W), and the laser light scanning speed S (unit: unit:). It is confirmed that it is preferable to control the relationship with mm / s) so as to satisfy "15 <h <150" and "67 (P / S) -3.5 <h <2222 (P / S) +13". .. That is, the hatched area is the pass determination area.

The above-described embodiments and experimental examples have been described for the purpose of assisting the understanding of the present invention, and the present invention is not limited to the specific configurations described. For example, it is possible to replace a part of the configuration of the embodiment with the configuration of the common general technical knowledge of those skilled in the art, and it is also possible to add the configuration of the common general technical knowledge of the person skilled in the art to the configuration of the embodiment. That is, the present invention may delete, replace with another configuration, or add another configuration to a part of the configurations of the embodiments and experimental examples of the present specification without departing from the technical idea of the invention. It is possible.

100 ... Turbine stationary blade,
101 ・ ・ ・ Inner ring side end wall, 102 ・ ・ ・ Wing part, 103 ・ ・ ・ Outer ring side end wall,
200 ・ ・ ・ Gas turbine, 210 ・ ・ ・ Compressor part, 220 ・ ・ ・ Turbine part,
221 ・ ・ ・ Turbine nozzle.

Claims (6)

A cobalt-based alloy material for metal lamination molding
With carbon of 0.08% by mass or more and 0.25% by mass or less,
Boron of 0.1% by mass or less and
Contains 10% by mass or more and 30% by mass or less of chromium
It contains 5% by mass or less of iron and 30% by mass or less of nickel, and the total of the iron and the nickel is 30% by mass or less.
It contains tungsten and / or molybdenum, and the total of the tungsten and the molybdenum is 5% by mass or more and 12% by mass or less.
It contains titanium, zirconium, niobium and tantalum, and the total of the titanium, the zirconium, the niobium and the tantalum is 0.5% by mass or more and 2% by mass or less.
With 0.5% by mass or less of silicon,
With 0.5% by mass or less of manganese,
Contains 0.003% by mass or more and 0.04% by mass or less of nitrogen.
The balance has a chemical composition of cobalt and impurities,
The impurities include 0.04% by mass or less of oxygen.
A cobalt-based alloy material characterized by that. In the cobalt-based alloy material according to claim 1,
The chemical composition is
The titanium content is 0.01% by mass or more and 1% by mass or less.
The zirconium is 0.05% by mass or more and 1.5% by mass or less.
The niobium is 0.02% by mass or more and 1% by mass or less.
The tantalum is 0.05% by mass or more and 1.5% by mass or less.
A cobalt-based alloy material characterized by that. In the cobalt-based alloy material according to claim 1 or 2.
The chemical composition further contains 0.5% by mass or less of aluminum as the impurities.
A cobalt-based alloy material characterized by that. The cobalt-based alloy material according to any one of claims 1 to 3.
The chemical composition further comprises 2% by weight or less of rhenium.
The total of the tungsten, the molybdenum and the rhenium is 5% by mass or more and 12% by mass or less.
A cobalt-based alloy material characterized by that. The cobalt-based alloy material according to any one of claims 1 to 4.
The cobalt-based alloy material, cobalt-based alloy material, characterized in that an alloy powder for selection択的laser melting method. In the cobalt-based alloy material according to claim 5,
A cobalt-based alloy material having a particle size of 5 μm or more and 100 μm or less of the alloy powder.

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