Nickel-cobalt-based alloy, nickel-cobalt-based alloy member using same, and method for manufacturing same

Abstract

The present invention provides an alloy that has excellent oxidation resistance and structural stability, high strength, and significantly improved service temperatures, and that is suitable for turbine disc uses. This nickel-cobalt-based alloy is characterized by having a composition comprising 15-43 mass% inclusive of cobalt, 6 mass% to less than 12 mass% of chromium, 3-9 mass% inclusive of tungsten, 1-6 mass% inclusive of aluminum, 1-8 mass% inclusive of titanium, 7 mass% or less of tantalum, 0.01-0.15 mass% inclusive of carbon, 0.01-0.15 mass% inclusive of boron, and 0.01-0.15 mass% inclusive of zirconium, with the remainder comprising nickel and unavoidable impurities.

Description

Nickel-cobalt-based alloy, nickel-cobalt-based alloy member using the same, and method for manufacturing the same

[0001] The present invention relates to a nickel-cobalt-based alloy, a nickel-cobalt-based alloy member using the alloy, and a method for producing the same.

[0002] Nickel-based alloys have been used for heat-resistant components, particularly turbine disks, of aircraft engines, power-generating gas turbines, and the like. Heat-resistant components such as turbine disks are required to have excellent strength, including creep strength and fatigue strength, as well as high-temperature oxidation resistance. Therefore, nickel-based alloys that are endowed with high-temperature oxidation resistance by adding chromium have been proposed. Known examples of the nickel-based alloy include, relative to the total amount, 11.5 to 11.9 mass% of Cr, 25 to 29 mass% of Co, 3.4 to 3.7 mass% of Mo, 1.9 to 2.1 mass% of W, 3.9 to 4.4 mass% of Ti, 2.9 to 3.2 mass% of Al, 0.02 to 0.03 mass% of C, 0.01 to 0.03 mass% of B, 0.04 to 0.06 mass% of Zr, 2.1 to 2.2 mass% of Ta, 0.3 to 0.4 mass% of Hf, 0.5 to 0.8 mass% of Nb, and the balance being Ni and unavoidable impurities (see Patent Document 1).

[0003] Further, as the nickel-based alloy, one containing, relative to the total amount, 20 to 40 mass% Co, 10 to 15 mass% Cr, 3 to 6 mass% Mo, 0 to 5 mass% W, 3.4 to 5 mass% Ti, 2.5 to 4 mass% Al, 0.01 to 0.05 mass% C, 0.01 to 0.05 mass% B, 0 to 0.1 mass% Zr, 1.35 to 2.5 mass% Ta, 0.5 to 1 mass% Hf, and 0 to 2 mass% Nb is known (see Patent Document 2).

[0004] Furthermore, as the nickel-based alloy, one containing, relative to the total amount, 11 to 15 mass% of Cr, 14 to 23 mass% of Co, 2.7 to 5 mass% of Mo, 0.5 to 3 mass% of W, 3 to 6 mass% of Ti, 2 to 5 mass% of Al, 0.015 to 0.1 mass% of C, 0.015 to 0.045 mass% of B, 0.015 to 0.15 mass% of Zr, 0.5 to 4 mass% of Ta, 0 to 2 mass% of Hf, and 0.25 to 3 mass% of Nb is also known (see Patent Document 3).

[0005] WO 2012 / 063879 U.S. Patent Application Publication No. 2009 / 0087338 European Patent Application Publication No. 1195446

[0006] Hirofumi Harada and Michio Yamazaki, "Alloy design of gamma prime precipitation strengthened Ni-base heat-resistant cast alloys containing Ti, Ta, and W," Iron and Steel, Vol. 65, pp. 1059-1068 (1979); Hidehiro Onodera, Hoichi Ryo, Toshihiro Yamagata, and Michio Yamazaki, "Effect of strength and ductility on high-temperature low-cycle fatigue of cast Ni-base cast alloys," Iron and Steel, Vol. 71, pp. 85-91 (1985)

[0007] Since the 1970s, the present applicant has been developing conventionally cast (CC), directionally solidified (DS), and single crystal (SC) nickel-based alloys suitable for turbine blades. By effectively utilizing the database accumulated through these efforts and an independently developed "alloy design program," the applicant has successfully developed nickel-based alloys with excellent high-temperature properties for both CC and SC materials (see, for example, Japanese Patent No. 3814662). More recently, the applicant has successfully developed a heat-resistant component suitable for turbine disk applications, with excellent temperature resistance (a temperature providing a lifespan of 1000 hours at 630 MPa), using the nickel-based alloy described in Patent Document 1, which is manufactured by powder metallurgy. The developed powder metallurgy (PM) material and cast and wrought (C&W) material having the same composition have excellent service temperature characteristics as the PM material and C&W material, respectively.

[0008] In recent years, PM materials, which have a homogeneous structure, have become the primary choice for turbine disks in high-temperature, high-pressure applications, which require high reliability. This is because the rapid solidification process in the gas atomization powder manufacturing process allows for alloy design with high additions of heavy metals such as Mo, Hf, Nb, and Ta, which are strengthening elements whose addition amounts are limited in C&W materials due to segregation. However, the service temperature of PM materials has only improved by 22°C over 28 years, from 693°C (IN100) in 1974 to 715°C (ME3) in 2002 (an improvement rate of 0.79°C / year). Furthermore, the addition of Mo, Hf, Nb, Ta, and other elements can reduce structural stability and oxidation resistance. Therefore, further improvements in engine performance require the development of alloys based on new design methods.

[0009] The present inventors conceived the present invention based on the idea that the service temperature of a turbine disk alloy could be significantly improved by following an alloy design based on a nickel-based alloy (known as a TM alloy) developed by the applicant in the 1970s for turbine blades, rather than by improving existing turbine disk alloys as in the past. Specifically, turbine blade alloys are designed for use at temperatures above 900°C, and therefore tend to have better oxidation resistance and corrosion resistance than turbine disk alloys intended for use at temperatures below 700°C. Therefore, the inventors conceived the present invention based on the turbine blade alloy, and believed that if the oxidation resistance and structural stability could be improved by reducing the addition of Mo, Nb, and Hf, and if strength could be achieved by adding other elements, a turbine disk alloy with significantly improved service temperature could be realized.

[0010] [1] The nickel-cobalt-based alloy of the present invention has a composition of 15% by mass or more and 43% by mass or less of cobalt, 6% by mass or more and less than 12% by mass of chromium, 3% by mass or more and 9% by mass or less of tungsten, 1% by mass or more and 6% by mass or less of aluminum, 1% by mass or more and 8% by mass or less of titanium, 7% by mass or less of tantalum, 0.01% by mass or more and 0.15% by mass or less of carbon, 0.01% by mass or more and 0.15% by mass or less of boron, 0.01% by mass or more and 0.15% by mass or less of zirconium, the balance being nickel and unavoidable impurities.

[0011] [2] In the nickel-cobalt-based alloy [1] of the present invention, the composition preferably consists of 15% by mass or more and 43% by mass or less of cobalt, 6% by mass or more and less than 12% by mass of chromium, 3% by mass or more and 9% by mass or less of tungsten, 1% by mass or more and 6% by mass or less of aluminum, 1% by mass or more and 8% by mass or less of titanium, 1.7% by mass or more and 7% by mass or less of tantalum, 0.01% by mass or more and 0.15% by mass or less of carbon, 0.01% by mass or more and 0.15% by mass or less of boron, 0.01% by mass or more and 0.15% by mass or less of zirconium, and the remainder being nickel and unavoidable impurities.

[0012] [3] In the nickel-cobalt-based alloy [1] or [2] of the present invention, the composition preferably consists of 15% by mass or more and 35% by mass or less of cobalt, 7% by mass or more and less than 12% by mass of chromium, 5.5% by mass or more and 8.5% by mass or less of tungsten, 2% by mass or more and 4% by mass or less of aluminum, 4% by mass or more and 7% by mass or less of titanium, 1.7% by mass or more and 7% by mass or less of tantalum, 0.01% by mass or more and 0.15% by mass or less of carbon, 0.01% by mass or more and 0.15% by mass or less of boron, 0.01% by mass or more and 0.15% by mass or less of zirconium, and the remainder being nickel and unavoidable impurities.

[0013] [4] In the nickel-cobalt-based alloys [1] to [3] of the present invention, the composition preferably consists of 15% to 35% by weight of cobalt, 8% to less than 12% by weight of chromium, 6.0% to 8.0% by weight of tungsten, 2% to 4% by weight of aluminum, 4% to less than 6.1% by weight of titanium, 1.7% to 7% by weight of tantalum, 0.01% to 0.15% by weight of carbon, 0.01% to 0.15% by weight of boron, 0.01% to 0.15% by weight of zirconium, and the remainder being nickel and unavoidable impurities. This composition [4] is based on the composition ranges of Example Alloys 1 to 3 described below and includes the composition range of Base Alloy 1 + (5% to 30%) x (Co-12.5% ​​by weight of Ti).

[0014] [5] In the nickel-cobalt-based alloys [1] to [4] of the present invention, it is preferable that at least one selected from the group consisting of less than 1.5% by mass of molybdenum, 5% by mass or less of niobium, and 2% by mass or less of hafnium is further contained as an optional composition element.

[0015] [6] In the nickel-cobalt-based alloys [1] to [5] of the present invention, it is preferable that the alloy further contains, as an optional composition element, at least one selected from the group consisting of 0.5 mass % or less of vanadium, 0.1 mass % or less of silicon, 0.05 mass % or less of calcium, 0.2 mass % or less of yttrium, and 0.2 mass % or less of a lanthanoid element.

[0016] [7] Among the nickel-cobalt-based alloys [1] to [6], it is preferable that the precipitates harmful to strength do not precipitate at 1.0 vol % or more, preferably 0.5 vol % or less, and more preferably 0.1 vol % or less, in a structural stability test at 850°C for 3000 hours. Precipitates harmful to strength include TCP (Topologically Close Packed) phase, beta phase, and Laves phase. [8] Among the nickel-cobalt-based alloys [1] to [6], it is preferable that the precipitates harmful to strength do not precipitate at 800°C for 3000 hours, more preferably 0.5 vol % or less, and more preferably 0.1 vol % or less. Precipitates harmful to strength include TCP (Topologically Close Packed) phase, beta phase, and Laves phase. －５The 0.2% yield strength at 1000 kJ / s is preferably 780 MPa or more, more preferably 890 MPa or more, and even more preferably 900 MPa or more.

[0017] [9] The nickel-cobalt-based alloy member of the present invention is made of any one of the nickel-cobalt-based alloys [1] to [6].

[10] The nickel-cobalt-based alloy member [9] of the present invention may be a turbine disk.

[0018]

[11] A method for producing a nickel-cobalt-based alloy member of the present invention comprises producing a nickel-cobalt-based alloy member [9] by one or more methods of casting including directional solidification, ordinary forging, powder metallurgy, additive manufacturing using powder, and three-dimensional manufacturing using wire.

[12] A method for producing a nickel-cobalt-based alloy member of the present invention comprises producing a nickel-cobalt-based alloy turbine disk

[10] by one or more methods of casting including directional solidification, ordinary forging, powder metallurgy, additive manufacturing using powder, and three-dimensional manufacturing using wire.

[13] In the method for producing a nickel-cobalt-based alloy member of the present invention

[11] or

[12] , preferably, the metal powder is produced by gas atomization.

[0019] The reasons for limiting the composition elements and their contents of the nickel-cobalt-based alloy of the present invention as described above will be described below. In the following description, % representing the content is mass %.

[0020] Cobalt (Co) is a useful component for controlling the solvus temperature of the γ' (gamma prime) phase of nickel-based alloys and nickel-cobalt-based alloys. Increasing the amount of cobalt lowers the γ' solvus temperature, widening the process tolerance range and improving forgeability. In particular, when titanium is contained in large amounts, it is desirable to add a slightly larger amount of cobalt to suppress the TCP phase and improve high-temperature strength. However, since an excessively high cobalt content reduces high-temperature strength, the cobalt content is typically 15% by mass or more and 43% by mass or less, preferably 15% by mass or more and 35% by mass or less.

[0021] Chromium (Cr) is added to improve environmental resistance and fatigue crack propagation properties. A Cr content of less than 6% by mass does not provide the desired properties, while a Cr content of more than 12% by mass tends to produce precipitates that reduce strength. Therefore, the Cr content is set to 6% by mass or more and 12% by mass or less, preferably 7% by mass or more and 12% by mass or less, and more preferably 8% by mass or more and 12% by mass or less.

[0022] Tungsten (W) dissolves in the γ phase and the γ' phase, strengthening both phases and being an effective element for improving high-temperature strength. If the W content is less than 3 mass%, sufficient improvement in high-temperature strength cannot be obtained. On the other hand, if the W content exceeds 9 mass%, high-temperature corrosion resistance may be reduced. Therefore, the tungsten content is 3 mass% or more and 9 mass% or less, preferably 5.5 mass% or more and 8.5 mass% or less, and more preferably 6 mass% or more and 8 mass% or less.

[0023] Aluminum (Al) is an element that promotes the formation of the γ' phase, and the amount of the γ' phase is mainly adjusted by the aluminum content. The aluminum content is 1% by mass or more and 6% by mass or less, preferably 2% by mass or more and 4% by mass or less. In addition, since the titanium to aluminum content ratio is strongly related to the formation of the harmful η phase, it is preferable to increase the aluminum content as much as possible to suppress this. Furthermore, aluminum is a raw material for forming aluminum oxide on the surface of nickel-based heat-resistant superalloys, and also contributes to improving oxidation resistance.

[0024] Titanium (Ti) is a desirable additive element for strengthening the γ' phase and improving strength. Its combined addition with cobalt provides nickel-based alloys and nickel-cobalt-based alloys with excellent phase stability and high strength. Specifically, by adding a cobalt-titanium alloy (e.g., Co-12.5% ​​by mass Ti used in the examples described below) with a similar γ+γ' two-phase structure to a nickel-based alloy or nickel-cobalt-based alloy with a γ+γ' two-phase structure, the combined addition of cobalt and titanium can be efficiently carried out, resulting in nickel-based alloys and nickel-cobalt-based alloys with high titanium content, stable structures, and high strength. The titanium content is typically 1% by mass or more and 8% by mass or less, preferably 4% by mass or more and 7% by mass or less, and more preferably 4% by mass or more and 6.1% by mass or less.

[0025] Tantalum (Ta) mainly substitutes for Al sites in the γ' phase and contributes to precipitation strengthening. On the other hand, if the tantalum content exceeds 7 mass%, harmful phases such as σ phase and μ phase tend to form, resulting in a decrease in high-temperature strength. For this reason, the tantalum content is 7 mass% or less. If the tantalum content is less than 1.7 mass%, precipitation strengthening is often insufficient, so the tantalum content is preferably 1.7 mass% or more and 7 mass% or less.

[0026] Carbon (C) is an element effective in improving ductility and creep properties at high temperatures. Typically, the carbon content is 0.01% by mass or more and 0.15% by mass or less. Boron (B) can improve creep properties, fatigue properties, etc. at high temperatures. Typically, the boron content is 0.01% by mass or more and 0.15% by mass or less. Carbon and boron exceeding the above content ranges may reduce creep strength or narrow the process tolerance range.

[0027] Zirconium (Zr) is an element effective in improving ductility, fatigue properties, etc. The zirconium content is usually 0.01 mass % or more and 0.15 mass % or less.

[0028] Molybdenum (Mo) mainly strengthens the γ phase and improves creep properties. Because molybdenum is a dense element, excessive molybdenum content increases the density of nickel-based alloys and nickel-cobalt-based alloys, which is undesirable from a practical standpoint. Typically, the molybdenum content is less than 1.5 mass%.

[0029] Niobium (Nb) is an effective strengthening element that contributes to reducing the specific gravity, but if the content is too high, harmful phases may be formed or quench cracks may occur at high temperatures. Usually, the niobium content is 5 mass% or less.

[0030] Hafnium (Hf) is a grain boundary segregating element that segregates to the grain boundaries to strengthen the grain boundaries, thereby improving high-temperature strength. If the hafnium content exceeds 2 mass%, it may cause local melting and reduce high-temperature strength, which is undesirable. Therefore, the hafnium content is 2 mass% or less.

[0031] Vanadium (V) is an element that mainly dissolves in the γ' phase and strengthens the γ' phase. The V content is preferably 0.5 mass% or less. If the V content exceeds 0.5 mass%, the creep strength decreases.

[0032] Silicon (Si) forms SiO on the alloy surface. ２ The silicon content is 0.1% by mass or less. If the silicon content exceeds 0.1% by mass, the solid solubility limit of other elements is reduced, and the required thermo-mechanical fatigue (TMF) and creep properties cannot be obtained.

[0033] Calcium (Ca), yttrium (Y), and lanthanide elements all contribute to the formation of Cr on the alloy surface. ２ O ３ Protective coating and Al ２ O ３These elements improve the adhesion of the protective coating, particularly the oxidation resistance during repeated oxidation. Therefore, these elements may be added as needed. However, excessive contents of these elements increase the amount of inclusions such as oxides, resulting in reduced hot workability and weldability. Therefore, the calcium content is set to 0.05% or less, and the yttrium and lanthanoid element contents are set to 0.2% or less. The above V, Si, Ca, Y, and lanthanoid elements may be contained alone or in combination of two or more.

[0034] Use of the nickel-cobalt-based alloy of the present invention will enable the extension of the service life of components used in high-temperature environments, such as gas turbine components, and will also enable the improvement of the efficiency of the gas turbine body when used in more severe environments. It is also expected that the alloy will be applied to hydrogen blend combustion / hydrogen gas turbines, ammonia blend combustion gas turbines, etc., which are being planned for the future.

[0035] 1 is a diagram showing a comparison of the yield of powder with a particle size of 53 μm or less for example alloys and some comparative alloys. 2 is a diagram showing an example of the relative density of HIP materials of the prepared base alloys and example alloys. 3 is a photograph showing the appearance of PM extrusion samples of example alloys 2 and 5. 4 is an inverse pole figure orientation map (IPF map) showing a comparison of the microstructures of example alloys 2 and 5 before and after PM extrusion. 5 is a scanning electron microscope (SEM) photograph showing the microstructure of example alloy 2 (aged at 870°C). 6 is a scanning electron microscope (SEM) photograph showing the microstructure of example alloy 2 (two-step aging). 7 is a scanning electron microscope (SEM) photograph showing the microstructure of comparative alloy 2 (two-step aging). 8 is a scanning electron microscope (SEM) photograph showing the microstructure of example alloy 2 (two-step aging). 9 is a scanning electron microscope (SEM) photograph showing the results of structural observation of some example alloys aged for a long period of time (3000 hours), showing the cases of aging temperatures of 650, 750, and 850°C for base alloy 1 and example alloys 1 to 4, respectively. 1 is a scanning electron microscope (SEM) photograph showing the structural stability of some comparative alloys during long-term aging (materials aged for 3000 hours), showing the results for Comparative Alloy 4 and Comparative Alloy 5 at aging temperatures of 650, 750, and 850°C. 2 is a diagram showing the results of a 375 MPa creep rupture test at 800°C for the comparative alloys and the base alloy. 3 is a diagram showing the results of a 375 MPa creep rupture test at 800°C for each of Base Alloys 1 and 2 and Example Alloys 1 to 7. 4 is a scanning electron microscope (SEM) photograph showing the initial structure of Example Alloy 4. 5 is a diagram showing creep curves for C&W materials of Example Alloy 2 and Comparative Alloy 2, showing the entire curve up to the region where creep strain exceeds 5%. 6 is a diagram showing creep curves for C&W materials of Example Alloy 2 and Comparative Alloy 2, showing the curve with an enlarged area up to the region where creep strain is 1%. This figure shows the service temperatures of a C&W alloy manufactured by a conventional casting and forging process, a conventional nickel-cobalt-based alloy manufactured by a powder metallurgy process, and Example Alloy 2, with the service temperatures on the vertical axis and the year of development on the horizontal axis. This figure also shows the high-temperature compressive strength of the base alloy and the Example Alloy, showing 0.2% yield strength at 725°C and 800°C. This figure also shows the results of an isothermal oxidation test on some of the Example Alloys.

[0036] The nickel-cobalt-based alloy and nickel-cobalt-based alloy member of the present invention will be described in more detail below with reference to examples, but the present invention is not limited to the following examples.

[0037] <Design of Example Alloys and Explanation (Selection) of Comparative Alloys> The example alloys were designed to achieve the above-mentioned composite addition of cobalt and titanium by adding a cobalt-titanium alloy (Co-12.5 mass% Ti) to a nickel-based alloy (base alloy) having a γ+γ' two-phase structure. Candidate base alloys were selected from the TM alloy group developed by the applicant for turbine blades. In this case, from the viewpoint of the ability to form grain boundary strengthening carbides and the disk weight, the alloys were selected to have a Cr content of 9.5 mass% or more, which is the main carbide forming element, and a density of 8.5 g / cm. ３ The selection criteria were as follows. Table 1 shows the compositions of base alloy 1 (see Non-Patent Document 1), selected from the TM alloy group, and a group of alloys (Example alloys 1 to 5) in which Co-12.5% ​​by mass Ti was added to base alloy 1 (see Non-Patent Document 2), as well as a group of alloys (Example alloys 6 and 7) in which Co-12.5% ​​by mass Ti was added to base alloy 2 (see Non-Patent Document 2). For example, the composition of Example alloy 1 is obtained by adding 10% Co-12.5% ​​by mass Ti to the composition of base alloy 1 (see the comments section of Table 1), and its Co content (17.3% by mass) is obtained by adding 87.5% by mass (Co content of the above cobalt-titanium alloy) x 0.1 = 8.75% by mass to 9.5% by mass (Co content of base alloy 1) x 0.9 = 8.55% by mass. Similarly, the Ti content of Example Alloy 1 is calculated as (3.90 × 0.9) + (12.5 × 0.1) = 4.76. The contents of Cr, W, Al, and Ta are each obtained by multiplying the content in Base Alloy 1 by 0.9. Note that Example Alloy 5 has a composition obtained by adding 20% ​​Co-12.5% ​​Ti by mass to the composition of Base Alloy 1 (i.e., the composition of Example Alloy 2), and further adding 0.7% Al by mass. Furthermore, to verify the superiority of the Example Alloys, commercially available Comparative Alloys 1 and 2, which are C&W materials, and Comparative Alloys 3, 4, and 5, which are PM materials, were selected as comparative alloys. The compositions of these alloys are also shown in Table 1.

[0038] <Composition list of example alloys and comparative alloys (single crystal materials)>

[0039] <Production of Example Alloys and Comparative Alloys (Single Crystal Materials)> For comparative evaluation, SC materials with compositions excluding the grain boundary strengthening elements C, B, and Zr from each alloy composition were cast in a unidirectional solidification furnace to eliminate the effects of the manufacturing process and structure. The cast single crystal test material round bars (10 mm diameter, 130 mm length) were subjected to solution treatment at 1200-1300°C for 5 hours, followed by aging heat treatment at 870°C for 20 hours. After heat treatment, test specimens for tensile creep tests (comparative evaluation tests) were taken from each single crystal test material round bar.

[0040] <Prototype Production of Example Alloys and Comparative Alloys (Sintered Powder Materials)> To evaluate the suitability of the example alloys and comparative alloys for the PM material manufacturing process, powders were prepared using a confined gas atomizer for some of the example alloys and some of the comparative alloys. The powders with particle sizes of 53 μm or less were classified, sealed, and then isotropically pressed (HIP) at 1100°C to obtain sintered bodies. Figure 1 shows the yield of powders with particle sizes of 53 μm or less used for sintering. For all example alloys, fine powders with particle sizes of 53 μm or less were produced, with a yield exceeding 70%. Furthermore, fine powders were obtained at a higher yield than the commercial alloy, Comparative Alloy 4. It was also confirmed that the yield of the powders increased with increasing amounts of Co-12.5% ​​Ti added to the base alloy. In the production of powder sintered materials, the higher the yield of powder used for sintering, the higher the production yield, so the example alloys are evaluated as alloy systems that are more suitable for powder production (low-cost and mass production) than the comparative alloys. Figure 2 shows an example of the relative densities of the HIP materials of the produced base alloy and example alloy. It can be seen that good high-density sintered bodies were obtained for both base alloy 1 and example alloy 2 shown in Figure 2.

[0041] <Prototypes of Example Alloys and Comparative Alloys (Forged Materials After Powder Sintering)> Among the example alloys prototyped as powder sintered materials, HIP sintered bodies of Example Alloy 2 and Example Alloy 5 were covered with carbon steel (SS400) and stainless steel (SUS304) to prepare extrusion billet specimens. These extrusion billet specimens were extruded using existing extrusion equipment in Japan, and were successfully extruded without clogging at an area reduction rate Re = 80% ( Figure 3 ). Figure 4 shows inverse pole figure (IPF) orientation maps showing the microstructures of Example Alloy 2 and Example Alloy 5 before (after heat treatment) and after extrusion. The "average grain size" in the figure refers to the average grain size. Data acquisition using the EBSD (Electron Backscattered Diffraction Pattern) method is carried out by AMETEK, Inc. EDAX Division, Model TEAM. ＴＭ The analysis was performed using EDS. The IPF analysis software was OIM version 8.1.0, manufactured by TSL Solutions Co., Ltd. According to the IPF map in Figure 4, both Example Alloy 2 and Example Alloy 5 had relatively uniform fine grains formed after extrusion, and no internal cracks were observed. It was also confirmed that PPB (Prior Particle Boundary), which causes a decrease in strength, had been eliminated. As described above, it was confirmed that the Example Alloys have good manufacturability using a powder metallurgy process.

[0042] <Prototype Production of Example Alloy and Comparative Alloy (Cast and Forged Materials)> Process technology development was carried out on Example Alloy 2 using vacuum melting, homogenization heat treatment, groove rolling, swaging, solution heat treatment, and aging heat treatment. For comparison, Comparative Alloy 2, a current C&W material, was also produced. Plastic processing was performed under preheating conditions at 1100°C to produce swaged materials. The swaged materials of Example Alloy 2 and Comparative Alloy 2 were subjected to solution heat treatments at 1160°C and 1100°C / 4h / AC, respectively. Furthermore, they were subjected to a two-step aging heat treatment of 650°C / 24h / AC + 760°C / 16h / AC. For comparison, a sample of Example Alloy 2 was also produced that was heat-treated at 870°C / 20h, the standard aging condition for single crystal model alloys. Figures 5A, 5B, and 5C show SEM photographs of the microstructures of each alloy. Figure 5A shows Example Alloy 2 (aged at 870°C), Figure 5B shows Example Alloy 2 (two-step aging), and Figure 5C shows Comparative Alloy 2 (two-step aging). SEM photographs were taken using a Carl Zeiss Gemini 300. As shown in Figure 5B, Example Alloy 2, which underwent two-step aging, exhibited typical fine grains pinned by primary γ' and a large amount of secondary γ' precipitates within the grains. As described above, various process conditions were investigated, and optimal process conditions were established that ultimately enabled the formation of a high-quality microstructure with a uniform grain size and a fine γ' precipitate distribution.

[0043] <Evaluation of Microstructural Stability (Sintered Powder Materials)> The solution-aged materials of the example alloys and comparative alloys, which were fabricated as sintered powder materials, were subjected to long-term aging tests at 650°C, 750°C, and 850°C, which are the expected operating temperatures for turbine disks. FIG. 6 shows scanning electron microscope (SEM) photographs of some of the example alloys aged for a long time (3000 hours), showing the results of aging at 650°C, 750°C, and 850°C for Base Alloy 1 and Example Alloys 1-4, respectively. FIG. 7 shows scanning electron microscope (SEM) photographs of some of the comparative alloys aged for a long time (3000 hours), showing the results of aging at 650°C, 750°C, and 850°C for Comparative Alloys 4 and 5, respectively. Table 2 also shows the results of measurements of precipitates (e.g., TCP phase, beta phase, Laves phase) that are harmful to strength, which precipitated during the microstructural stability test at 850°C for 3000 hours.

[0044] <Amount of harmful precipitates>

[0045] These results indicate that the structural stability of the powder sintered materials improves in the order of Example Alloy > Comparative Alloy. This is due to the amount of heavy elements, such as Mo, Nb, Ta, and Hf, which induce the precipitation of harmful third phases. In particular, the addition of Nb, Ta, and Hf, which are strengthening elements actively added to PM alloys manufactured by major manufacturers around the world, induces the precipitation of large amounts of harmful phases within and at grain boundaries, as can be seen from the structural observation results of Comparative Alloy 5 and Comparative Alloy 4 in Figure 7. On the other hand, the Example Alloy, which does not contain Mo, Nb, or Hf and contains less Ta than the current material, is concluded to be an extremely promising alloy system. Furthermore, in the base alloy, precipitation of TCP and beta phases was confirmed in Base Alloy 1 aged at 850°C. In contrast, in Example Alloys 1 to 5, in which Co-12.5 mass% Ti was added to Base Alloy 1, the Co and Ti contents were increased compared to Base Alloy 1, resulting in a relative decrease in the Ta and W contents, and thus improved structural stability. From the above, it can be concluded that Example Alloys 1, 2, 3, and 4 in particular exhibit excellent structural stability in the operating temperature range for turbine disk applications. From the above, it can be concluded that the nickel-cobalt-based alloy of the present invention should have the property that, in a structural stability test at 850°C for 3,000 hours, precipitates harmful to strength do not precipitate in 1.0 vol% or more, preferably 0.5 vol% or less, and more preferably 0.1 vol% or less.

[0046] <Creep properties of example alloys and comparative alloys (single crystal material)> Using prototype single crystal materials, tensile creep rupture tests were conducted at 800°C under a load stress of 735 MPa. Figure 8 shows the results of the tensile creep tests on the base alloy and comparative alloys. Under these test conditions, comparative alloy 3, an alloy developed by the applicant, showed the longest life among the comparative alloys, with a rupture life of t ｆ= 11.9 h. Furthermore, Comparative Alloy 1 had the shortest lifespan, with a rupture lifespan of 0.42 h. The order of rupture lifespan was Comparative Alloy 3 > Comparative Alloy 5 > Comparative Alloy 2 > Comparative Alloy 4 > Comparative Alloy 1, and from this result it can be seen that Comparative Alloy 3 and Comparative Alloy 2, which are nickel-cobalt based alloys, exhibit excellent creep properties. On the other hand, Base Alloy 1 and Base Alloy 2 exhibited extremely excellent creep properties compared to the comparative alloys. The rupture lives of Base Alloy 1 and Base Alloy 2 were t ｆ = 38.39 h and t ｆ 9A shows the results of a creep rupture test at 800°C under 375 MPa for each of base alloys 1 and 2 and example alloys 1 to 7. Focusing on the creep properties of base alloy 1 and example alloys 1, 2, 3, and 4, which are base alloy 1 with Co-12.5% ​​by mass Ti added, the creep life increased as the amount of Co-12.5% ​​by mass Ti added to base alloy 1 increased, and example alloy 3, which contains 30% by mass or more of Co, showed a t ｆ It was revealed that the maximum creep life was 123.22 h. Therefore, compared to conventional nickel-based or nickel-cobalt-based alloys for turbine blades, the combined addition of Co and Ti by adding a cobalt-titanium alloy (Co-12.5% ​​by mass Ti) with a similar γ+γ' dual-phase structure was proven to be extremely effective in improving creep life. On the other hand, Example Alloy 4, which contained 40% by mass Co, exhibited a shorter life than Base Alloy 1. This is due to the presence of a large amount of TCP phase (harmful phase), 2.1% by volume, in the final solidification portion of the initial structure, as shown in Figure 9B. Example Alloy 5, which added Co-12.5% ​​by mass Ti and Al to Base Alloy 1, exhibited a shortened life due to the remaining eutectic γ' phase. Furthermore, Example Alloys 6 and 7, which added Co-12.5% ​​by mass Ti to Base Alloy 2, also tended to exhibit a longer life as the amount of Co-12.5% ​​by mass Ti increased.

[0047] <Creep Properties (Cast and Forged Materials) of Example Alloys and Comparative Alloys> Figure 10A shows the creep properties of cast and forged materials of Example Alloy 2 and Comparative Alloy 2, showing the overall curve up to the region where creep strain exceeds 5%. Figure 10B shows an enlarged curve of the region up to 1% creep strain in Figure 10A. Figure 10A shows the creep rupture test results at 725°C and 630 MPa, with the creep time leading to creep rupture as the reference temperature. The creep rupture test results for Comparative Alloy 2 are 698°C, while the creep rupture test results for Example Alloy 2 aged at 870°C are 728°C, and the creep rupture test results for Example Alloy 2 aged at two stages are 739°C. Figure 10B shows the creep rupture test results at 725°C and 630 MPa, with the creep time leading to creep rupture as the reference temperature. The creep temperature limit for Comparative Alloy 2 is 727°C, while that for Example Alloy 2 aged at 870°C is 740°C, and that for Example Alloy 2 two-stage aged is 766°C. As shown in Figure 10A, Example Alloy 2 has a rupture life approximately 10 times longer than that of Comparative Alloy 2, which currently has the world's highest rupture temperature limit for a C&W material. Furthermore, the plastic working process enabled the decomposition of harmful PPB, resulting in a creep rupture elongation of approximately 5%. Furthermore, when comparing Example Alloy 2 with the other alloys, as shown in Figures 10A and 10B, the two-stage aged material, which exhibits finer precipitates, had better creep properties, whether evaluated by rupture life or 0.2% creep time.

[0048] <Service Temperatures of Example Alloys and Comparative Alloys (Cast and Forged Materials)> Figure 11 shows the service temperatures of a C&W alloy manufactured by a conventional cast and forged process, a conventional nickel-cobalt-based alloy manufactured by a powder metallurgy process, and Example Alloy 2. The vertical axis indicates the service temperature, and the horizontal axis indicates the year of development. In Figure 11, open circles (◯) indicate C&W alloys, and filled circles (●) indicate powder metallurgy-processed alloys. Figure 11 also shows the service temperatures of Example Alloy 2 and Comparative Alloy 2, which were prototyped as cast and forged materials. The creep life of Example Alloy 2 at 725°C / 630 MPa was approximately 2100 hours, which corresponds to a service temperature of 739°C in terms of creep rupture life. Furthermore, the service temperature, calculated based on a 0.2% creep life, which is important for disk alloys (turbine disk alloys) that are life-limited parts, was 766°C. Thus, Example Alloy 2 exhibited the world's highest service temperature rating for a cast and forged material at this stage. As mentioned above, the service temperature rating of PM materials increased by only 22°C over 28 years, from 693°C (IN100) to 715°C (ME3) between 1974 and 2002, representing an improvement rate of 0.79°C / year. Example Alloy 2's service temperature rating was improved by 24°C compared to ME3. From this, it can be concluded that the design guidelines for the new nickel-cobalt-based alloy, which uses a turbine blade alloy as the base alloy, are extremely effective as a technology for significantly improving the service temperature rating while maintaining the manufacturability of gas turbine disks.

[0049] <Compressive strength of example alloys and base alloys (sintered powder materials)> Compression tests were conducted on some of the base alloys and example alloys manufactured as sintered powder materials at temperatures of 725°C and 800°C. As shown in Figure 12, it was revealed that example alloys 1 to 5, which are base alloy 1 to which Co-12.5 mass% Ti was added, all exhibited a 0.2% yield strength higher than that of base alloy 1, except for example alloy 4. Here, the 0.2% yield strength was determined by the strain rate of 10 at the specified temperature. －５( / s), the 0.2% yield strength refers to the stress value at the stage where the strain reaches 0.2%. For base alloy 1, the 0.2% yield strengths at 725°C and 800°C were 980 MPa and 780 MPa, respectively. In contrast, for example alloy 1, the 0.2% yield strengths at 725°C and 800°C were 1100 MPa and 930 MPa, respectively. For example alloy 2, the 0.2% yield strengths at 725°C and 800°C were 1080 MPa and 890 MPa, respectively. For example alloy 3, the 0.2% yield strengths at 725°C and 800°C were 990 MPa and 795 MPa, respectively. For example alloy 4, the 0.2% yield strengths at 725°C and 800°C were 950 MPa and 750 MPa, respectively. In Example Alloy 5, the 0.2% yield strength at 800°C is 880 MPa, and no data is available for 725°C.

[0050] That is, the 0.2% proof stress of these example alloys increased significantly with increasing amounts of Co-12.5% ​​by mass Ti added to the base alloy, and then gradually decreased. It is believed that the alloy exhibiting the maximum strength in the design guidelines of this example based on base alloy 1 exists between example alloy 1 and example alloy 2. The increase in 0.2% proof stress up to example alloy 1 is believed to be the effect of increasing the precipitate phase volume fraction and Ti content. On the other hand, the gradual decrease from example alloy 2 to example alloy 4 is believed to be due to the addition of Co-12.5% ​​by mass Ti to base alloy 1, which reduces the contents of Cr and W, which are major solid solution strengthening elements in the alloy matrix.

[0051] <Oxidation Resistance of Example Alloys and Comparative Alloys (Sintered Powder Materials)> Solution-aged materials of the example alloys and comparative alloys, which were prototyped as sintered powder materials, were subjected to 500-hour isothermal oxidation tests at 650°C, 750°C, and 850°C. Figure 13 shows the weight gain results for example alloys 1, 2, 3, and 4. Adding Co-12.5% ​​Ti by mass to base alloy 1 reduces oxidation resistance. This reduction is believed to be due to a decrease in the content of Al and Cr, which form oxide films. However, the oxidation resistance of these example alloys compares favorably with that of comparative alloy 4, a commercial alloy, and is therefore not considered to pose a practical problem.

[0052] <Summary of Comparison Between Example Alloys and Comparative Alloys> As described above, the example alloys exhibit superior creep strength, service temperature, and compressive strength compared to the comparative alloys, and also have sufficient oxidation resistance. This confirms that the present inventors have made it possible to improve alloys for turbine disks based on a new design methodology based on alloys for turbine blades.

[0053] Although several embodiments of the present invention have been described, these embodiments are presented as examples and are not intended to limit the scope of the invention. These novel embodiments can be embodied in various other forms, and various omissions, substitutions, and modifications can be made without departing from the spirit of the invention. These embodiments and their modifications are included within the scope and spirit of the invention, and are also included in the scope of the invention and its equivalents as defined in the claims.

[0054] Based on the design guideline of adding a cobalt-titanium alloy (Co-12.5% ​​by mass Ti in the above-described example) with a similar γ+γ' dual-phase structure to a nickel-based alloy for turbine blades as the base alloy, the combined addition of cobalt and titanium has been shown to produce a novel nickel-cobalt-based alloy with excellent creep strength, service temperature, compressive strength, sufficient oxidation resistance, and good forgeability. Specifically, the above-described example alloys were shown to be more suitable for powder production (low-cost and mass production) than the comparative alloys. Furthermore, it was demonstrated that billets with a homogeneous microstructure can be produced by either the casting process or the powder metallurgy process, confirming industrial-level manufacturability. The use of the nickel-cobalt-based alloy of the present invention can extend the service life of components used in high-temperature environments, such as gas turbine components, and can improve the efficiency of gas turbine bodies and other components used in harsher environments. It is also expected that the alloy will be applicable to hydrogen-mixed combustion / hydrogen gas turbines, ammonia-mixed combustion gas turbines, and other applications planned for the future.

Claims

1. The composition is 15% by weight or more and 43% by weight or less of cobalt, 6% by weight or more and less than 12% by weight of chromium; 3% by mass or more and 9% by mass or less of tungsten; 1% by weight or more and 6% by weight or less of aluminum, 1% by mass or more and 8% by mass or less of titanium; 7% by weight or less of tantalum, 0.01% by mass or more and 0.15% by mass or less of carbon; 0.01% by mass or more and 0.15% by mass or less of boron; 0.01% by mass or more and 0.15% by mass or less of zirconium, A nickel-cobalt based alloy, the balance of which consists of nickel and unavoidable impurities.

2. The composition is 15% by weight or more and 43% by weight or less of cobalt, 6% by weight or more and less than 12% by weight of chromium; 3% by mass or more and 9% by mass or less of tungsten; 1% by weight or more and 6% by weight or less of aluminum, 1% by mass or more and 8% by mass or less of titanium; 1.7% by weight or more and 7% by weight or less of tantalum, 0.01% by mass or more and 0.15% by mass or less of carbon; 0.01% by mass or more and 0.15% by mass or less of boron; 0.01% by mass or more and 0.15% by mass or less of zirconium, The nickel-cobalt based alloy of claim 1, characterized in that the balance consists of nickel and unavoidable impurities.

3. The composition is 15% by weight or more and 35% by weight or less of cobalt, 7% by weight or more and less than 12% by weight of chromium; 5.5% by mass or more and 8.5% by mass or less of tungsten; 2% by weight or more and 4% by weight or less of aluminum, 4% by weight or more and 7% by weight or less of titanium; 1.7% by weight or more and 7% by weight or less of tantalum, 0.01% by mass or more and 0.15% by mass or less of carbon; 0.01% by mass or more and 0.15% by mass or less of boron; 0.01% by mass or more and 0.15% by mass or less of zirconium, The nickel-cobalt based alloy of claim 1, characterized in that the balance consists of nickel and unavoidable impurities.

4. The composition is 15% by weight or more and 33% by weight or less of cobalt, 8.9% by weight or more and less than 12% by weight of chromium; 6.0% by mass or more and 8.3% by mass or less of tungsten; 2.5% by weight or more and 3.5% by weight or less of aluminum; 4.4% or more and less than 6.5% by weight of titanium; 1.8% by mass or more and 2.5% by mass or less of tantalum, 0.01% by mass or more and 0.15% by mass or less of carbon; 0.01% by mass or more and 0.15% by mass or less of boron; 0.01% by mass or more and 0.15% by mass or less of zirconium, The nickel-cobalt based alloy of claim 1, characterized in that the balance consists of nickel and unavoidable impurities.

5. Further optional composition elements include less than 1.5% by weight of molybdenum, 5% by weight or less of niobium, and 2% by weight or less of hafnium, The nickel-cobalt-based alloy according to claim 1, characterized in that it contains at least one selected from the group consisting of:

6. Further optional composition elements include 0.5% by weight or less of vanadium, 0.1% by weight or less of silicon, 0.05% by weight or less of calcium, 0.2% by weight or less of yttrium, and 0.2% by weight or less of a lanthanide element, The nickel-cobalt-based alloy according to claim 1, characterized in that it contains at least one selected from the group consisting of:

7. 7. The nickel-cobalt-based alloy according to claim 1, wherein in a structural stability test at 850° C. for 3,000 hours, precipitates harmful to strength do not precipitate in an amount of 1.0% by volume or more.

8. The nickel-cobalt-based alloy according to any one of claims 1 to 6, －５ ( / s) 0.2% yield strength of 780 MPa or more.

9. A nickel-cobalt-based alloy member made of the nickel-cobalt-based alloy according to any one of claims 1 to 6.

10. The nickel-cobalt based alloy component of claim 9, wherein the nickel-cobalt based alloy component is a turbine disk.

11. A method for producing the nickel-cobalt-based alloy member according to claim 9 by one or more methods including casting including a directional solidification method, ordinary forging, powder metallurgy, additive manufacturing using a metal powder, and three-dimensional manufacturing using a wire.

12. A method for producing a nickel-cobalt-based alloy member, comprising the steps of: casting including a directional solidification method; ordinary forging; powder metallurgy; additive manufacturing using a metal powder; and three-dimensional manufacturing using a wire, the nickel-cobalt-based alloy turbine disk according to claim 10 being produced by one or more methods.

13. The method for producing a nickel-cobalt based alloy member according to claim 11, wherein the metal powder is a metal powder produced by a gas atomization method.

14. The method for producing a nickel-cobalt based alloy part according to claim 12, wherein the metal powder is a metal powder produced by a gas atomization method.