JP2023143401A - METHOD OF MANUFACTURING Ni ALLOY MEMBER - Google Patents

Abstract

To provide a Ni alloy member that exhibits small elongation at the initial stage of a creep test.SOLUTION: Provided is a method of manufacturing Ni alloy member that includes: an additive manufacturing step of forming, by an additive manufacturing technique, a member having a mass% composition of Ni: 50-55%, Cr: 17.0-21.0%, Nb+Ta: 4.75-5.5%, Mo: 2.8-3.3%, Ti: 0.65-1.15%, Al: 0.20-0.80%, Co: 1.0% or less, Cu: 0.3% or less, C: 0.08% or less, Si: 0.35% or less, Mn: 0.35% or less, P: 0.015% or less, S: 0.015% or less, B: 0.006% or less, and balance: Fe and unavoidable impurities; a recrystallization treatment step of recrystallizing the formed member at a temperature of 1120°C or higher and 1250°C or lower; and a solution treatment step of subjecting the recrystallized member to solution treatment at a temperature of 925°C or higher and 1010°C or lower.SELECTED DRAWING: Figure 1

Description

The present invention relates to a method of manufacturing a member made of Inconel 718 alloy (Inconel is a registered trademark, hereinafter abbreviated as "718 alloy"), which is a type of Ni-based super heat-resistant alloy, by additive manufacturing and heat treatment.

718 alloy is known as one of the Ni-based super heat-resistant alloys. This alloy is a precipitation hardening type alloy. Precipitation hardening alloys are heated above the solubility curve to form a uniform solid solution of solute atoms, and then subjected to aging treatment to precipitate another phase within the material, so that the precipitated phase undergoes no transition. By impeding movement, plastic deformation of the material is suppressed. In 718 alloy, fine crystal grains of γ'' phase (Ni ₃ Nb) and γ' phase (Ni ₃ (Al, Ti)) precipitate in the γ phase, which is the parent phase, due to aging treatment, resulting in super heat resistance. Sexuality is expressed.

718 alloy members are formed by forging, casting, etc., and in recent years are often formed by additive manufacturing technology. However, 718 alloy parts formed using additive manufacturing technology have poor creep properties, and the creep properties differ greatly between the stacking direction and the direction perpendicular to it, so there have been no attempts to improve the creep properties by developing heat treatment methods. It has been.

Patent Document 1 discloses that a member made of a precipitation hardening type Ni-Cr based Ni-based alloy formed by additive manufacturing is heated at a temperature of (T _C -100)°C or higher, where the solid solution line temperature is T _C , ( It is described that solution treatment is carried out at a temperature in the temperature range of T _C -50)°C or lower. In the example, 718 alloy (T _C = 1260°C) is solution-treated at 1180°C, so that the creep rupture time is uniform regardless of the angle with respect to the modeling direction, and is comparable to that of cast material. It is stated that this can be improved. This result is said to be due to the coarsening of crystal grains caused by the solution treatment at high temperatures.

Japanese Patent Application Publication No. 2017-203195

According to the method described in Patent Document 1, creep characteristics are improved by subjecting a 718 alloy member formed by additive manufacturing to solution treatment at a high temperature. However, according to the experimental results of the present inventors, it was found that even in test specimens subjected to high-temperature solution treatment, the elongation from transition creep at the initial stage of the creep test to steady creep was large. Referring to Figure 13, the creep curves in the test are as follows: instantaneous strain consisting of plastic strain appears at the moment of loading, transition creep where the strain rate decreases as work hardening progresses, and strain rate where work hardening and tissue recovery are balanced and the strain rate increases. It is classified into three stages: steady creep, which becomes constant, and accelerated creep, where the strain rate accelerates and leads to rupture. Among these, when the strain rate from transition creep to steady creep is high, the 0.2% elongation time to reach the elastic region limit becomes shorter, so this is not suitable for applications such as turbine blades that must maintain precise dimensions and minute voids. , there is a problem that the time during which the product can be used in a safe area is shortened. In addition, depending on the application of the 718 alloy member, the member may be subjected to solution treatment and aging treatment after it has been used for a certain period of time, and the member may be recycled, repaired, and reused. It is necessary to suppress the growth of In this regard, the method described in Patent Document 1 has room for further improvement.

The present invention has been made in consideration of the above, and an object of the present invention is to provide a technique that can suppress elongation at the initial stage of a creep test of a 718 alloy member formed by additive manufacturing.

The method for manufacturing the Ni alloy member of the present invention is based on mass percentage: Ni: 50-55%, Cr: 17.0-21.0%, Nb+Ta: 4.75-5.5%, Mo: 2.8-55%. 3.3%, Ti: 0.65 to 1.15%, Al: 0.20 to 0.80%, Co: 1.0% or less, Cu: 0.3% or less, C: 0.08% or less , Si: 0.35% or less, Mn: 0.35% or less, P: 0.015% or less, S: 0.015% or less, B: 0.006% or less, balance: Fe and inevitable impurities. A layered manufacturing step in which a member having a composition is formed by additive manufacturing technology, a recrystallization treatment step in which the formed member is recrystallized at a temperature of 1120° C. or higher and 1250° C. or lower, and the recrystallized member is It has a solution treatment step of performing solution treatment at a temperature of 925°C or higher and 1010°C or lower.

According to the method for manufacturing a Ni alloy member of the present invention, the strain rate at the initial stage of the creep test can be suppressed by subjecting the obtained Ni alloy member to aging treatment under standard conditions.

It is a process flow diagram of the Ni alloy member manufacturing method of one embodiment. It is a figure showing the shape of the test piece used for a tensile test and a creep test. It is a figure which shows the result of 0.2% elongation time in a creep test. It is a figure showing the result of the rupture time in a creep test. It is a creep curve of Comparative Example 3. 3 shows creep curves of Comparative Example 13 and Example 1. 1 is an image of crystal grains obtained by electron backscatter diffraction (EBSD). 3 is a SEM image of creep fracture surfaces of Comparative Example 3, Comparative Example 13, and Example 1. 3 is a SEM image of creep fracture surfaces of Comparative Example 3, Comparative Example 13, and Example 1. 3 is a SEM image of creep fracture surfaces of Comparative Example 3, Comparative Example 13, and Example 1. 3 is a SEM image of the cut surfaces of Comparative Example 3, Comparative Example 13, and Example 1. 3 is a SEM image of the cut surfaces of Comparative Example 3, Comparative Example 13, and Example 1. FIG. 3 is a diagram for explaining a creep curve.

An embodiment of the method for manufacturing a Ni alloy member of the present invention will be described along the process flow shown in FIG.

The 718 alloy powder used in this embodiment has Ni: 50 to 55%, Cr: 17.0 to 21.0%, Nb+Ta: 4.75 to 5.5%, and Mo: 2.8 to 3% by mass. .3%, Ti: 0.65 to 1.15%, Al: 0.20 to 0.80%, Co: 1.0% or less, Cu: 0.3% or less, C: 0.08% or less, Composition of Si: 0.35% or less, Mn: 0.35% or less, P: 0.015% or less, S: 0.015% or less, B: 0.006% or less, balance: Fe and inevitable impurities. has. This alloy is called Inconel 718 (Inconel is a registered trademark) and is listed in the Uniform Numbering System for Alloys (UNS) by the American Society for Testing and Materials (ASTM) and Society of Automotive Engineers (SAE) as N07718 and in JIS G4901 and 4902. It is specified as NCF718.

Using the above alloy powder, a member is shaped by additive manufacturing technology. As the additive manufacturing method, preferably a laser layered manufacturing method (SLM method) is used. The SLM method is a type of powder bed fusion bonding method, in which the raw material alloy powder is spread on a modeling stage to form a uniform thin layer, and the thin layer is irradiated with scanning laser light to melt the alloy powder.・By repeating solidification, alloy layers are stacked to form a component.

A member formed by the SLM method contains many columnar crystal grains extending in the stacking direction of the alloy due to the forming method. Therefore, mechanical properties such as creep differ between the stacking direction and the direction perpendicular to the stacking direction. In the following, the stacking direction during modeling will be referred to as the Z direction, and the direction perpendicular to the stacking direction will be referred to as the XY direction. In addition, in the SLM method, in order to suppress the influence of deviation in the scanning direction of the laser beam, the scanning direction is rotated by a predetermined angle for each layer and lamination is performed, so the structure of the model is created in a plane perpendicular to the Z direction. So it is isotropic. Also in this specification, the XY direction only means perpendicular to the Z direction, and does not mean a specific direction within a plane perpendicular to the Z direction.

Next, the additively manufactured member (as-shaped material) is recrystallized. The purpose of the recrystallization process is to reduce the anisotropy in the Z and XY directions as much as possible. Specifically, the purpose is to reduce the difference in size of crystal grains in the Z direction and the XY direction by recrystallization. In the following, this may be referred to as isotropic or equiaxed.

The temperature of the recrystallization treatment is 1120°C or higher, preferably 1140°C or higher. This allows the member to be recrystallized. On the other hand, the temperature of the recrystallization treatment is 1250°C or lower, preferably 1200°C or lower, more preferably 1180°C or lower. If the recrystallization treatment temperature exceeds 1250° C., there is a risk that the member will soften and deform during the treatment. Moreover, if the recrystallization treatment temperature exceeds 1200° C., the crystal grains become too large and the rupture time in the creep test becomes short. When the recrystallization treatment temperature is in the range of 1140° C. to 1180° C., equiaxed crystal grains progress most.

In the recrystallization treatment, the time for maintaining the member at the above temperature is preferably 30 minutes or more, more preferably 1 hour or more. This is to allow sufficient time for recrystallization. On the other hand, in the recrystallization treatment, the time for maintaining the member at the above temperature is preferably 2 hours or less. Although the recrystallization treatment time does not have as great an effect as the temperature, if the recrystallization treatment time is too long, the crystals will become too large. The member is held in an inert gas atmosphere or a vacuum atmosphere at normal pressure.

After recrystallizing the member by holding it at a predetermined temperature for a predetermined time, the member is cooled to at least about 600° C. or lower by blowing air or inert gas onto the member. Moreover, it is preferable that the cooling rate is 10° C./min or more. This makes it possible to prevent the δ phase and the like from precipitating during the cooling process. As described later, the presence of an appropriate amount of δ phase is considered to be useful for improving creep properties, but since it is technically difficult to control the amount of δ phase precipitated by the cooling rate, recrystallization treatment Preferably, no phase is precipitated.

Next, the recrystallized member is subjected to solution treatment. The purpose of solution treatment is to uniformly dissolve solute atoms into a solid solution. Furthermore, in the solution treatment of 718 alloy, it is common to precipitate fine crystal grains of the δ phase (Ni ₃ Nb). The δ phase is said to be stable at temperatures below about 1010-1020°C. Since the δ phase has the same constituent elements as the γ'' phase, it is said that if the δ phase grows too much, the amount of γ'' phase precipitated during subsequent aging treatment will be reduced and the effect of precipitation hardening will be negated. ing. On the other hand, the δ phase is said to precipitate at the grain boundaries of the parent phase and suppress the growth of crystal grains, and is thought to have the function of improving creep characteristics.

The temperature of the solution treatment is preferably 925°C or higher and 1010°C or lower. This temperature is a standard solution treatment temperature for 718 alloy parts formed by forging or the like. For example, according to SAE's ASM5662 standard, parts are held at 1725-1850F (F is Fahrenheit, equivalent to 940-1010°C) and then air cooled. Furthermore, the reference appendices of JIS G4901 and 4902 describe conditions for treatment at 925 to 1010° C. and rapid cooling as solution heat treatment. This makes it possible to uniformly solutionize the structure of the member and precipitate an appropriate amount of δ phase.

In the solution treatment, the time for maintaining the member at the above temperature is preferably 30 minutes or more, more preferably 1 hour or more. This is to allow sufficient time for solution treatment. On the other hand, in the solution treatment, the time for maintaining the member at the above temperature is preferably 2 hours or less. Although the solution treatment time does not have as great an effect as the temperature, if the time is too long, the δ phase will grow too much.

After the member is held at a predetermined temperature for a predetermined time to form a solution, the member is cooled by blowing air or inert gas onto the member. This is to maintain the solution state by cooling it somewhat rapidly.

Next, the recrystallized member is subjected to an aging treatment. The purpose of the aging treatment is to precipitate the γ'' phase and the γ' phase in the γ phase, which is the parent phase.

The temperature of the aging treatment is preferably 610°C or higher and 730°C or lower. This temperature is a standard aging treatment temperature for 718 alloy members formed by forging or the like. For example, according to SAE's ASM5663 standard, a member is held at 1325F (718°C) for 8 hours, lowered to 1150F (621°C) at 100F/h, held at that temperature for another 8 hours, and air cooled. In addition, in the reference appendix of JIS G4901 and 4902 mentioned above, after the solution heat treatment, it is maintained at 705 to 730°C for 8 hours, furnace cooled to 610 to 630°C, and then aged at that temperature and then air cooled for a total aging time. It is stated that the period is 18 hours.

In addition, according to the SAE ASM5664 standard, as a solution treatment, the component is held at 1900 to 1950F (1038 to 1066℃) and then air cooled, and as an aging treatment, the component is measured at 1400F (760℃) and 1200F (649℃). Although a 20-hour treatment method is specified, this standard is a rather special standard, as it is intended to be used at low temperatures at the expense of creep characteristics at high temperatures.

The Ni alloy member of this embodiment is obtained through the above layered manufacturing, recrystallization treatment, solution treatment, and aging treatment.

Using an alloy powder having the composition shown in Table 1 as a raw material, a member was modeled using a powder additive manufacturing system (EOS GmbH, M290) using a Yb fiber laser (focal diameter: 80 μm). For the Z-direction tensile test and creep test, the test piece shown in FIG. 2 was directly shaped with the length direction as the lamination direction. For tensile tests in the XY directions and creep tests, test pieces shown in FIG. 2 were cut out from the shaped rectangular parallelepiped.

The samples of Examples 1 and 2 were prepared by changing the recrystallization temperature and subjecting them to recrystallization treatment, solution treatment, and aging treatment. In the recrystallization treatment, the sample was held at a predetermined temperature in a vacuum atmosphere for 1 hour, and then cooled to room temperature by air cooling. In the solution treatment, the sample was held at 980° C. for 1 hour in a vacuum atmosphere, and then cooled to room temperature by air cooling. The aging treatment was performed by holding at 720°C for 8 hours in a vacuum atmosphere, furnace cooling to 620°C for 2 hours, holding at 620°C for a further 8 hours, and cooling to room temperature by air cooling.

The sample of Comparative Example 1 is an as-shaped material that has not been subjected to heat treatment. Samples of Comparative Examples 2 to 16 were prepared by solution treatment and aging treatment at different solution temperatures. In the solution treatment, the sample was held at a predetermined temperature in a vacuum atmosphere for 1 hour, and then cooled to room temperature by air cooling. The aging treatment was performed in the same manner as in the examples.

The tensile test was conducted using an autograph in accordance with ASTM E21 at a temperature of 650° C. and a strain rate of 0.005/s. The creep test was conducted under the conditions of 621° C. and a load of 724 MPa in accordance with JIS Z2271.

Table 2 shows the results of the tensile test at 650°C.

Naturally, the tensile strength of Comparative Example 1 (as-shaped material) was lower than the other samples in both the 0.2% proof stress and the tensile strength. Comparing the results of Comparative Examples 2 to 16, it is found that as the solution treatment temperature is raised from 950°C, the strength increases in both the Z direction and the It reached its highest at ℃ and gradually decreased above that temperature. Further, when the solution temperature was 1100°C or lower, strength was higher in the XY direction than in the Z direction, and at 1120°C or higher, the difference between the Z direction and the XY direction was reduced. The cause of the change in strength due to the difference in solution treatment temperature in Comparative Examples 2 to 16 is not clear, but the as-built material contains the δ phase and the Laves phase (Fe ₂ (Nb, Ti)), which is brittle and easily becomes a crack initiation point. It is possible that these phases were dissolved by heat treatment at 1020°C or higher. Moreover, when the solution treatment temperature exceeded 1030° C., there is a possibility that the tensile strength gradually decreased as the crystal grains became larger. The 0.2% proof stress and tensile strength of Examples 1 and 2 showed almost the same values as those of the comparative example, which was solution treated at the same temperature as the recrystallization treatment.

Table 3 shows the results of the creep test at 621°C. Further, among the results shown in Table 3, FIG. 3 shows the results of 0.2% elongation time, and FIG. 4 shows the results of rupture time. The horizontal axes in FIGS. 3 and 4 are the recrystallization treatment temperature for Examples and the solution treatment temperature for Comparative Examples. Further, FIG. 5 shows the creep curves of Comparative Example 3, and FIG. 6 shows the creep curves of Comparative Example 13 and Example 1.

From Table 3, FIG. 3, and FIG. 4, in the comparative example, the creep characteristics in the XY direction are extremely poor when the solution temperature is 1100° C. or lower, and the difference with the Z direction is narrowed when the solution temperature is 1120° C. or higher. When comparing Examples and Comparative Examples when the recrystallization temperature of the Examples and the solution treatment temperature of the Comparative Examples are similar, there is a large variation in the data, but the rupture times (Figure 4) are almost the same. , 0.2% elongation time (FIG. 3) tended to be longer in Examples. The 0.2% elongation time of a sample of 718 alloy forged material solution-treated and aged under standard conditions is said to be about 200 to 400 hours, and the 0.2% elongation time of Examples 1 and 2 was It has been improved to the same extent.

Comparative Example 3 in FIG. 5 (solution temperature 980° C. + aging treatment) was heat treated by the method described in the ASM standard and JIS standard, but the creep performance was poor. It is shown that solution annealing and aging of additively manufactured 718 alloy shapes under standard conditions does not result in good creep properties. In modeling using additive manufacturing, Nb segregation occurs at the cell interface of the dendrite structure that is formed when a melt pool containing molten alloy powder is rapidly solidified, and the δ phase grows too much under standard solution treatment and aging conditions. it is conceivable that.

Comparing Comparative Example 13 (solution temperature 1160°C + aging treatment) and Example 1 (recrystallization temperature 1160°C + solution temperature 980°C + aging treatment) in Fig. 6, the rupture time in both the Z direction and the XY direction is was longer in Comparative Example 13, but the strain rate from transition creep at the initial stage of the test to steady creep was lower in Example 1. This is also reflected in Table 3, where Example 1 has a longer 0.2% elongation time than Comparative Example 13.

The above results confirm that recrystallization treatment at high temperature and standard solution treatment for 718 alloy suppress elongation at the initial stage of the creep test, rather than one-time solution treatment at high temperature. did it.

In order to investigate the mechanism that brought about the above results, the crystal structure of the sample was investigated.

First, in order to investigate the effect of heat treatment temperature on the shape and size of crystals, samples were prepared in which the as-shaped material was held at 980 to 1200°C for 1 hour and air-cooled, and the as-built material was held at 1160°C for 1 hour and air-cooled. After that, the sample was held at 980° C. for 1 hour and cooled in air, and then backscattered electron diffraction (EBSD) measurement was performed on a cross section parallel to the Z direction. According to EBSD, it is possible to know the crystal orientation of each crystal grain, and it is possible to create an image showing the outline of the crystal grain on the measurement surface. FIG. 7 is an image showing grain boundaries. FIG. 7 shows a grain boundary where the crystal orientations on both sides differ by 10 degrees or more. The vertical direction in FIG. 7 is the Z direction of the sample, the horizontal direction is the XY direction of the sample, and the length of the bar at the bottom left of each image is 600 μm. Table 4 shows the average values of the widths in the Z direction and the XY direction of the crystal grain regions appearing on the measurement surface. Since the measurement plane divides each crystal grain into two at random positions, the average value of this width does not indicate the average value of the size of each crystal grain in the Z direction or the XY direction, but it does serve as an index. . Note that each sample in FIG. 7 and Table 4 has not been subjected to aging treatment, and by aging treatment, it becomes a sample of each comparative example or example shown in the remarks column.

From FIG. 7 and Table 4, the crystal grains become larger as the heat treatment temperature increases, and the difference in size between the Z direction and the XY direction becomes smaller at 1100 to 1200° C., especially at 1120 to 1180° C. Further, when comparing the crystal sizes of those heat-treated at 1160°C and those heat-treated at 1160°C + 980°C, there was not much difference, although the latter was slightly larger. Considering this result together with the results of the tensile test and creep test of Examples and Comparative Examples, when the values in Table 4 are approximately 40 μm, that is, the heat treatment temperature is 1140 to 1180°C, the tensile properties and creep properties are isotropic. It is considered to be the most effective and to show the best results.

Next, the fracture surfaces of the creep test pieces of Comparative Example 3, Comparative Example 13, and Example 1 shown in FIGS. 5 and 6 were observed. Figures 8 to 10 show scanning electron microscope (SEM) photographs of the fractured surface. In FIGS. 8 to 10, Z means a creep test in which a tensile load was applied in the Z direction, and XY means a creep test in which a tensile load was applied in the XY directions. The length of the bar under each SEM image is 1 mm in FIG. 8, 300 μm in FIG. 9, and 100 μm in FIG. 10.

From FIGS. 8 to 10, there was no fracture originating from a clear defect on any of the fracture surfaces, and it is presumed that the difference in the shape of the fracture surfaces is due to the difference in crystal structure and plastic deformation. Except for Comparative Example 3 Z, fracture at grain boundaries was dominant in all test specimens. In addition, in the SEM photograph, the fracture surface caused by intergranular fracture shows irregularities due to the shape of the crystal grains, while the fracture surface caused by intragranular fracture appears to be a rough, flat surface with minute dimples formed. . Grain boundary fracture occurred almost entirely in XY of Comparative Example 3, and in Z of Comparative Example 3, there were many locations where intragranular fracture occurred (for example, the central part of Comparative Example 3Z in FIG. 10). In Comparative Example 13, grain boundary fracture occurred throughout both Z and XY. In Example 1, grain boundary fracture occurred mostly in both Z and XY, but there were a few areas where intragranular fracture occurred (for example, the lower left of Example 1Z in FIG. 10 and the center of Example 1XY in the same figure). .

11 and 12 show SEM photographs of the cut surfaces of Comparative Example 3, Comparative Example 13, and Example 1. The cut plane is a plane parallel to the Z direction, and the vertical direction in the figure is the Z direction.

In FIG. 11, bright points in the image are the δ phase or Laves phase. In Comparative Example 3 (solution temperature: 980° C. + aging treatment), a large number of δ phases are present along the grain boundaries. Comparative Example 13 (solution temperature 1160°C + aging treatment) and Example 1 (recrystallization temperature 1160°C + solution temperature 980°C + aging treatment) also have a δ phase along the grain boundaries, but in Comparative Example 3 Fewer. Furthermore, in the X-ray diffraction (XRD) measurements of Comparative Example 13 and Example 1, very small peaks of metal carbides were observed, suggesting that carbides such as Cr, Ti, and Nb were also precipitated at grain boundaries. .

FIG. 12 is an even higher resolution SEM photograph. From FIG. 12, the structure composed of the precipitated phases of γ' phase and γ'' phase becomes finer in the order of Comparative Example 3, Comparative Example 13, and Example 1. The reason for this is presumed as follows. In the as-built material, there is segregation of solute elements due to the cellular or dendrite-like structure.In Comparative Example 3, the precipitated phase containing the solute elements is unevenly distributed due to the strong influence of the structure of the as-built material.Comparative Example 13 In Example 1, the composition becomes more uniform due to the heat treatment at 1160°C, so the precipitated phase becomes finer.Furthermore, in Example 1, there is less segregation during cooling after the solution treatment at 980°C, and A fine precipitated phase is obtained.

From the above, it is thought that the following changes in the structure occur during the manufacturing process of the Ni alloy member of the embodiment. By recrystallizing the as-shaped material at a temperature of 1120° C. or higher, equiaxing in the Z direction and the XY direction progresses to some extent, and the anisotropy of the tensile properties and creep properties is reduced. If the recrystallization treatment temperature is 1140 to 1180°C, equiaxed formation will proceed more. Moreover, during this recrystallization treatment at high temperature, the δ phase and the Laves phase are dissolved, and the segregation of solute elements progresses to become uniform. In the recrystallized member, fine δ phase is precipitated by solution treatment. As a result, the strain rate at the initial stage of the creep test is suppressed by precipitating fine γ'' and γ' phases during the aging treatment.

The present invention is not limited to the above-described embodiments and examples, and various modifications can be made within the scope of the technical idea.

Claims (2)

In mass%, Ni: 50 to 55%, Cr: 17.0 to 21.0%, Nb + Ta: 4.75 to 5.5%, Mo: 2.8 to 3.3%, Ti: 0.65 to 1.15%, Al: 0.20 to 0.80%, Co: 1.0% or less, Cu: 0.3% or less, C: 0.08% or less, Si: 0.35% or less, Mn: A member having a composition of 0.35% or less, P: 0.015% or less, S: 0.015% or less, B: 0.006% or less, and the remainder: Fe and unavoidable impurities is formed using additive manufacturing technology. Additive manufacturing process;
a recrystallization treatment step of recrystallizing the shaped member at a temperature of 1120° C. or higher and 1250° C. or lower;
a solution treatment step of solutionizing the recrystallized member at a temperature of 925° C. or higher and 1010° C. or lower;
A method for manufacturing a Ni alloy member having the following. further comprising an aging treatment step of heat treating the solution-treated member at a temperature of 610° C. or higher and 730° C. or lower;
The method for manufacturing a Ni alloy member according to claim 1.

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