Method for producing precipitation hardening austenitic alloy steel material, hot-worked precipitation hardening austenitic alloy material, and precipitation hardening austenitic alloy steel material

Abstract

The present invention provides a precipitation hardening austenitic alloy steel material which is capable of further improving mechanical characteristics without improving the composition of a steel material, and with which it is possible to expect a longer service life of a component in a high-temperature high-load environment or a high-pressure hydrogen environment.　Provided is a method for producing a precipitation hardening austenitic alloy steel material, the method comprising: a hot working step in which hot working is performed on a material for hot working so as to obtain a hot-worked material that has crystal grains having a GOS value of 1.0° or more at an area ratio of 50% or more, the material having a composition of a precipitation hardening austenitic alloy; and a heat treatment step in which the hot-worked material is subjected to solution heat treatment and aging treatment so as to obtain a steel material. The solution heat treatment temperature in the heat treatment step is 850-1050°C, and the aging treatment temperature in the heat treatment step is 700-750°C. Also provided are: a hot-worked precipitation hardening austenitic alloy material; and a precipitation hardening austenitic alloy steel material.

Description

Manufacturing method of precipitation hardened austenitic alloy steel, hot-worked precipitation hardened austenitic alloy material, and precipitation hardened austenitic alloy steel

[0001] The present invention relates to a method for producing a precipitation hardened austenitic alloy steel, a hot-worked precipitation hardened austenitic alloy material, and a precipitation hardened austenitic alloy steel.

[0002] Precipitation-hardened austenitic alloys such as SUH660 are known as components suitable for applications such as bolts and shafts due to their excellent strength characteristics over a wide temperature range. They are also known as components suitable for hydrogen stations (hydrogen stands) due to their excellent resistance to hydrogen embrittlement in hydrogen environments. For example, Patent Document 1 discloses that A286 alloy (equivalent to SUH660), a γ'-phase precipitation-hardened alloy, is suitable as an alloy with high mechanical properties related to resistance to hydrogen embrittlement in hydrogen or when hydrogen is absorbed. Patent Document 2 also discloses that A286 alloy (equivalent to SUH660), suitable for hydrogen energy equipment, is forged at a total forging ratio of 5:1 to obtain a forged product.

[0003] As described in the aforementioned patent documents, precipitation-hardened austenitic alloys have excellent resistance to hydrogen embrittlement. However, in high-temperature environments or in high-pressure hydrogen gas, the presence of hydrogen causes localized deformation due to plastic deformation, which can easily lead to the formation of stacking faults. Such defects are likely to cause cracks and tend to reduce mechanical properties such as tensile strength compared to those in air. For example, Non-Patent Document 1 describes that when precipitation-hardened A286 test specimens were hydrogen-charged, the tensile strength and other properties of the test specimens decreased with increasing hydrogen content. Given this background, further strengthening of precipitation-hardened austenitic alloys is also required. Non-Patent Document 2 describes improving the composition of the A286 alloy to suppress the precipitation of the η phase, which increases susceptibility to hydrogen embrittlement.

[0004] Patent Document 1: JP 2011-068919 A, Chinese Patent Application Publication No. 11354982

[0005] Naoki Tajima and five others, "Effect of Internal Hydrogen on Tensile Properties of Iron-Based Superalloy SUH660," Transactions of the Japan Society of Mechanical Engineers (Series A), August 2012, Vol. 78, No. 792, pp. 50-64; Shinya Sato and three others, "Development of Improved A286 Alloy with Excellent High-Temperature Strength and Evaluation of Hydrogen Embrittlement Susceptibility," Japan Steel Works Technical Report, Japan Steel Works, Ltd., December 2014, No. 65, pp. 76-81

[0006] The method described in Non-Patent Document 2 is effective in improving high-temperature mechanical properties, but there is a concern that the increased strength may result in a decrease in ductility, such as elongation. Furthermore, Non-Patent Document 2 contains a relatively large amount of W to improve properties. However, W is an expensive raw material, and its use should be avoided as much as possible in industrial products. Furthermore, in order to meet the needs of applications such as jet engine turbine rotors, heat-resistant bolts, and automotive heat-resistant parts, not only higher strength but also improved creep rupture strength, which indicates the ability to resist deformation in high-temperature environments, is required. Therefore, an object of the present invention is to provide a method for producing a precipitation-hardened austenitic alloy steel that can further improve mechanical properties without modifying the steel composition and that can be expected to extend the life of parts in harsh high-temperature, high-load environments and high-pressure hydrogen environments, making it suitable for use in aircraft, energy, and automotive applications.

[0007] The present invention has been made in view of the above-mentioned problems. That is, one aspect of the present invention is a method for producing a precipitation hardened austenitic alloy steel material, comprising: a hot working step of hot working a hot working material having a composition of a precipitation hardened austenitic alloy to obtain a hot worked material having crystal grains with an area ratio of 50% or more having a GOS value of 1.0° or more; and a heat treatment step of solution treating the hot worked material and aging treating it to obtain a steel material, wherein the solution treatment temperature in the heat treatment step is 850 to 1050°C and the aging treatment temperature in the heat treatment step is 700 to 750°C. Another aspect of the present invention is a hot-worked precipitation hardened austenitic alloy material having an area ratio of 50% or more of crystal grains with a GOS value of 1.0° or more, and having a 0.2% yield strength of 590 MPa or more, a rupture life of 30 hours or more, and a gamma prime phase mean circle equivalent diameter of 10 to 35 nm when solution treated at 850 to 1050°C and aging treated at a temperature of 700 to 750°C. Another aspect of the present invention is a precipitation hardened austenitic alloy steel material having a 0.2% yield strength of 590 MPa or more, a rupture life of 30 hours or more, and a gamma prime phase mean circle equivalent diameter of 10 to 35 nm.

[0008] According to the present invention, it is possible to produce a precipitation hardened austenitic alloy steel material that can further improve mechanical properties without modifying the composition of the steel material and that is expected to extend the life of parts under high-temperature, high-load environments and high-pressure hydrogen environments.

[0009] It is a schematic diagram showing the sample collection position for EBSD in the example. It is a schematic diagram showing the sample collection position for mechanical property measurement in the example. It is a schematic diagram for explaining the wrought forming ratio. It is a schematic diagram for explaining other wrought forming ratios.

[0010] The present invention will be described in detail below. However, the present invention is not limited to the embodiments described herein, and appropriate combinations and improvements are possible within the scope of the technical concept of the invention. The present invention is directed to precipitation-hardened austenitic alloy steel. The precipitation-hardened austenitic alloy steel refers to SUH309, SUH310, SUH330, SUH660, SUH661, and their improved versions as described in JIS-G-4311. SUH309, SUH310, SUH330, SUH660, and SUH661 are preferred. SUH660 is more preferred. Specifically, the steel preferably contains, by mass, 10 to 40% Ni and 10 to 30% Cr, with the total of Fe, Ni, and Cr being 95% or more by mass. Alternatively, the alloy may contain 10 to 40% Ni and 10 to 30% Cr, with the remainder being Fe and unavoidable impurities. A more preferred lower limit for the Ni content is 20% by mass, and a more preferred upper limit for the Ni content is 30% by mass. A more preferred upper limit for Cr is 20% by mass. Furthermore, in order to improve hardness, high-temperature strength, and the like, one or more elements of V, Si, Mn, Al, B, Ta, W, Ti, Mo, and Nb may be contained in a total amount of up to 5.0% by mass. Other unavoidably contained impurity elements include, for example, C, S, P, and O, and the upper limit for each of these elements is preferably set to 0.1%. Preferably, the alloy contains 10 to 40% Ni, 10 to 30% Cr, and contains one or more elements of V, Si, Mn, Al, B, Ta, W, Ti, Mo, and Nb in a total amount of up to 5.0% by mass, with the remainder being Fe and unavoidable impurities.

[0011] In the manufacturing method of this embodiment, a material for hot working having a precipitation-hardened austenitic alloy composition is first prepared. This material for hot working is preferably a steel ingot obtainable by casting. Here, hot working in this embodiment refers to hot plastic working of the steel ingot, such as hot pressing or hot rolling, followed by machining into a round bar or square bar. Note that if the machined steel slab (billet, bloom, etc.) is subjected to homogenization heat treatment, this steel slab may be used as a material for forging. Furthermore, when producing a steel ingot, remelting may be performed to reduce elemental segregation and nonmetallic inclusions.

[0012] The prepared hot-worked material is then heated in a heating furnace and hot-worked multiple times to obtain a hot-worked material. The hot-worked material of the present invention is characterized by having 50% or more crystal grains with an area ratio of 1.0° or more (GOS (Grain Orientation Spread) value). A preferred area ratio is 60% or more, and a more preferred area ratio is 70% or more. This GOS value can be measured using the conventionally known "SEM-EBSD method (Electron Backscatter Diffraction)" and can be derived by calculating the misorientation of the points (pixels) that make up the crystal grains. The crystal misorientation obtained by the GOS value is an index of the strain imparted to the alloy by processing. When 50% or more crystal grains have an area ratio of 1.0° or more, a uniform fine-grained structure with little variation in crystal grain size tends to be obtained after solution treatment. Furthermore, there are fewer coarse unrecrystallized grains that have not been subjected to processing strain, which tends to improve mechanical properties, particularly 0.2% proof stress. On the other hand, if the area ratio of crystal grains with a GOS value of 1.0° or more is less than 50%, the above-mentioned coarse unrecrystallized grains increase, which may result in a decrease in mechanical properties such as 0.2% proof stress. The GOS value in this embodiment can be derived by taking a sample at a depth of D / 4 or D / 2 from the surface of the hot-worked material in the axial direction (D is the distance between the surfaces passing through the axis; for example, it corresponds to the diameter in the case of a cylindrical shape and the diagonal length in the case of a rectangular column) and measuring it. The cross sections for observing the area ratio include cross sections perpendicular to the axis and cross sections in the axial direction. However, it is preferable that the area ratio of crystal grains with a GOS value of 1.0° or more is 50% or more when observed in both the cross sections perpendicular to the axis and the axial cross sections of the bar. Furthermore, the area ratio of crystal grains having a GOS value of 1.0° or more after solution heat treatment and aging heat treatment tends to be a low value of less than 50% because strain due to hot working is released. Note that the GOS value described above in this embodiment can be measured by taking a test piece from the surface of the obtained hot-worked material at a depth D / 4 or D / 2 in the axial direction in the direction shown in Figure 1. In particular, the depth D / 2 corresponds to the center of the steel material and is the least susceptible to strain, so it is preferable that the test piece taken from this depth D / 2 position also satisfies the characteristics specified in the present invention.

[0013] In order to obtain the GOS value of the present invention described above, it is preferable to perform hot working so that the total hot working ratio during hot working is 15 or more. This allows working strain to be introduced into the center of the hot-worked material, and recrystallization occurs due to dynamic recrystallization, thereby improving the mechanical properties of the hot-worked material. Representative examples of "hot working" in this embodiment include hot forging and hot rolling, and hot forging and hot rolling can also be combined. When hot forging is selected, the total hot working ratio of the present invention indicates the total forging ratio (hereinafter also referred to as the total forging ratio). The "total forging ratio (total forging ratio)" in the present invention is derived using the forging ratio of solid forging described in JIS-G-0701. For example, when the hot forging shown in FIG. 13 is performed, ((S ０ / S １ ) × (S １ / S ２ )×・・・・・×(S ｎ－１ / S ｎ )) is the total forging ratio. ０ , S １ , S ２ ....S ｎ－１ , S ｎ is the cross-sectional area. As an example of a process in which the total forging ratio exceeds 15, when a 100 mm diameter round bar is used as the forging material, if substantial forging is performed from 100 mm to 50 mm, and then from 50 mm to 24 mm, the total forging ratio is 17.3, which satisfies the requirements of the present invention. There is no particular upper limit to the total forging ratio, but since the greater the number of forgings, the higher the manufacturing cost, a realistic setting of 200 is practical, and 150 is also possible. When hot rolling is selected for hot working, the cross-sectional areas before and after rolling can be measured and applied to the formula for deriving the forging ratio for substantial forging described above to determine the hot working ratio. Here, since hot rolling can increase the working ratio at a lower cost than hot forging, the total hot working ratio when hot rolling is applied can be increased compared to when hot forging is applied. (For example, the total hot working ratio is 1000, etc.) Hereinafter, in this embodiment, unless otherwise specified, hot forging is selected as the hot working.

[0014] In this embodiment, when hot forging is applied to the hot working, it is preferable to perform upset forging at least once in order to increase the total forging ratio without drastically reducing the size of the forged material. In this embodiment, since the processing strain imparted by substantial forging tends to primarily contribute to improving the mechanical properties, when upset forging is performed, the forging ratio (upset ratio) during upset forging is not included in the calculation of the total forging ratio in this embodiment. For example, when hot forging as shown in FIG. 4 is performed (S １ -S ２ Between the two is upset forging), ((S ０ / S １ ) × (S ２ / S ３ )×・・・・・×(S ｎ－１ / S ｎ )) is the total forging ratio. By combining upset forging, sufficient processing strain can be introduced even for large forged products with an equivalent circle diameter of 100 mm or more. The temperature of the heating furnace during hot working is preferably 800 to 1300°C, and the surface temperature of the material at the end of hot working is preferably 600°C or higher.

[0015] In the present invention, "hot forging" preferably includes hot free forging, in which a workpiece is placed on an anvil with a flat or curved surface and is then pressed or struck with an anvil or die, also with a flat or curved surface, to form the material. Of course, when forming complex shapes, die forging using a die may be performed, or radial forging, in which the workpiece is rotated in the circumferential direction while being pressed from four directions along its entire length to obtain a forged material, may be performed. A combination of these processes may also be used. Furthermore, to achieve the total forging ratio described above, it is preferable to perform solid forging at least three times. This is because precipitation-hardened austenitic alloy steel has high deformation resistance, so a smaller reduction rate per forging allows for more stable processing. Furthermore, it is preferable to perform upset forging at least twice to introduce processing strain into the interior of large forged products, even those with an equivalent circle diameter of 100 mm or more.

[0016] In this embodiment, the hot-worked material obtained by the hot working step is subjected to solution treatment and aging treatment to obtain a precipitation-hardened austenitic alloy steel material. First, the solution treatment in this embodiment is performed at 850 to 1050°C. The solution treatment temperature may be adjusted to suit the desired properties. When emphasis is placed on yield strength, the solution treatment temperature is preferably 850 to 950°C, more preferably 870 to 930°C. When emphasis is placed on creep rupture strength, the solution treatment temperature is preferably 950 to 1050°C, more preferably 970 to 1030°C.

[0017] The aging treatment temperature in this embodiment is 700 to 750°C. By carrying out the aging treatment within the narrow temperature range described in the present invention after carrying out the above-mentioned hot working step and solution treatment, it is possible to improve the tensile properties, particularly the yield strength, as much as possible while also appropriately increasing the creep rupture strength. A preferred lower limit of the aging treatment temperature is 710°C, and a more preferred lower limit of the aging treatment temperature is 715°C. A preferred upper limit of the aging treatment temperature is 745°C.

[0018] The precipitation-hardened austenitic alloy steel material of the present invention can be a steel material having a columnar shape and an equivalent-area circle diameter of 10 mm or more in a cross section perpendicular to the axis. The steel material of the present invention obtained by the manufacturing method of the present invention has an equivalent-area circle diameter of 10 mm or more, and can be applied to materials of various sizes. Furthermore, since the manufacturing method of the present invention can introduce processing strain to the center of a large steel material, the equivalent-area circle diameter is preferably 25 mm or more. More preferably, it is 50 mm or more, 100 mm or more, 200 mm or more, 300 mm or more, or 400 mm or more. Applying the manufacturing method of the present invention to a steel material having an equivalent-area circle diameter of 100 mm or more is particularly preferable, since it is possible to introduce processing strain to the center of the material even in large steel materials. Here, the term "columnar shape" refers to, for example, a cylindrical or rectangular column shape, and may be tapered in the axial direction. When a tapered shape is formed, the equivalent-area circle diameter of the cross section perpendicular to the axis of the present invention can be measured at a position L / 2, where L is the length of the steel material.

[0019] The precipitation hardened austenitic alloy steel material of the present invention obtained by the manufacturing method of the present invention has a 0.2% yield strength of 590 MPa or more. When it is desired to improve creep rupture strength as well, the precipitation hardened austenitic alloy steel material obtained by the manufacturing method of the present invention preferably has a 0.2% yield strength of 590 to 630 MPa. A preferable lower limit of the 0.2% yield strength is 610 MPa. Furthermore, if a minimum creep rupture strength (for example, a rupture strength of 30 hours or more, as described below) is sufficient, the precipitation hardened austenitic alloy steel material obtained by the manufacturing method of the present invention can have a 0.2% yield strength of 655 MPa or more (preferably 670 MPa or more). Furthermore, the elongation is preferably 20% or more, more preferably 23% or more, and even more preferably 25% or more. Furthermore, the tensile strength is preferably 900 MPa or more, and more preferably 950 MPa or more. Furthermore, the reduction of area is preferably 20% or more, and more preferably 30% or more.

[0020] The precipitation-hardened austenitic alloy steel material of the present invention obtained by the manufacturing method of the present invention preferably has a grain size number of 4.0 or more. The upper limit of the grain size number is not particularly limited, but in order to obtain excessively fine grains, the hot working ratio must also be set excessively high, which makes manufacturing difficult when hot forging is mainly used. Therefore, the upper limit may be set to, for example, 10.0. A preferred upper limit of the grain size number is 8.0, and a more preferred upper limit of the grain size number is 7.0. In particular, the 0.2% yield strength is affected by the grain size and tends to decrease due to the presence of coarse unrecrystallized grains that have not been given sufficient processing strain. However, the steel material of the present invention obtained by the manufacturing method of the present invention can achieve a good value without decreasing the 0.2% yield strength. The 0.2% yield strength, elongation, tensile strength, and grain size described above in this embodiment, as with the GOS value, are measured by taking test specimens from the surface of the obtained forged material at a depth of D / 4 or D / 2 in the axial direction in the direction shown in FIG. 2 .

[0021] The precipitation hardened austenitic alloy steel material of the present invention has a high yield strength as described above, and also exhibits a good creep rupture strength. This creep rupture strength is the rupture life (time to rupture) obtained by performing a creep rupture test according to ASTM E139. The precipitation hardened austenitic alloy steel material of the present invention has an excellent rupture life in the creep rupture test, such as a rupture life of 30 hours or more. This rupture life indicates the rupture life when the creep rupture test is performed under the conditions of a test temperature of 649°C and a load stress of 448 MPa, and a test temperature of 649°C and a load stress of 483 MPa. The precipitation hardened austenitic alloy steel material obtained by the manufacturing method of the present invention maintains a high yield strength as described above, and also exhibits a good rupture life. Preferably, the rupture life is 50 hours or more, and more preferably, the rupture life is 80 hours or more. The longer the fracture life, the better. However, if the fracture life is extremely long, tensile properties such as yield strength tend to decrease, so the upper limit may be set to, for example, 300 hours. A preferred upper limit is 290 hours. A more preferred upper limit is 280 hours.

[0022] In the precipitation-hardened austenitic alloy steel of the present invention, the gamma prime phase has an equivalent circle diameter of 10 nm to 35 nm. In particular, in order to exhibit good mechanical properties, the average equivalent circle diameter of the gamma prime phase is preferably 11.0 nm or more, and preferably 15.0 nm or more. Furthermore, the average equivalent circle diameter of the gamma prime phase is preferably 34.0 nm or less. When the equivalent circle diameter of the gamma prime phase is less than 10 nm, the 0.2% proof stress tends to be insufficient depending on the working ratio. This is presumably because the gamma prime phase does not sufficiently contribute to the prevention of dislocation glide within the steel. Furthermore, when the equivalent circle diameter of the gamma prime phase exceeds 35 nm, the fracture time may be significantly reduced depending on the heat treatment conditions. This is presumably because dislocation glide becomes easier due to a change in the mechanism of dislocation glide within the steel.

[0023] The present invention will be described in more detail in the following examples. (Example 1) A steel ingot (equivalent to SUH660) having the composition shown in Table 1 was prepared as a raw material for hot-worked columns for use in examples of the present invention. Solid forging and upset forging were performed multiple times so that the total hot working ratio was 15.8, resulting in a hot-worked columnar material having an equivalent diameter of a circle of 160 mm in the cross section perpendicular to the axis of the columnar material. Two types of samples were then taken from the surface of the hot-worked material (before heat treatment) at depths D / 4 and D / 2 in the axial direction, one in the axial direction and the other in the perpendicular direction, and the GOS value of each sample was confirmed. The GOS value was measured using a ZEISS field emission scanning electron microscope and a TSL OIM (Orientation-Imaging-Micrograph) EBSD measurement and analysis system. The longitudinal section (axial section) and transverse section (cross section perpendicular to the axis) of the sample were observed, and measurement samples (11 × 10 × 5 mm) were taken from each cross section, as shown in Figure 1. The measurement surface was 11 × 10 mm, the measurement field of view was 1500 μm × 1500 μm, and the step distance between adjacent pixels was 3.0 μm. Observation was performed under conditions that determined whether a boundary with an orientation difference of 5° or more between adjacent pixels was a grain boundary. From the resulting GOS value map, the area ratio of crystal grains with a GOS value of 1.0° or more relative to the entire observation field was calculated. Table 2 shows the observation results. The results in Table 2 show that the percentage of the obtained hot-worked material with a GOS value of 1.0° or more is 50% or more in both the longitudinal section and the transverse section.

[0024] Next, test specimens (15 mm diameter x 100 mm) were taken from the surface of the resulting hot-worked material in the direction shown in Figure 2 ("Material taken from the axial direction") at a depth of D / 4 (D: diameter of the circle equivalent to the area diameter) in the axial direction. They were solution-treated and aged, and then various mechanical properties and grain size numbers were measured. Test specimens with a circular cross section as specified in ASTM E8 were used to measure 0.2% proof stress, tensile strength, elongation, and reduction of area, and tests were conducted in accordance with ASTM E8's Tensile Test Method for Metallic Materials. Test specimens with a circular cross section as specified in ASTM E139 were used to measure rupture time (rupture life), and tests were conducted in accordance with ASTM E139's Creep Rupture Test Method for Metallic Materials. The grain size number was determined using the grain size standard plate I in accordance with JIS-G-0551. The average equivalent circle diameter of the gamma prime phase was determined by observing the specimens at 100,000 or 200,000 magnifications using a scanning electron microscope (SEM) and randomly selecting 10 specimens from within the field of view. Table 3 shows the mechanical properties and grain size numbers. The specimens solution-treated at 899°C were designed with an emphasis on yield strength, and creep rupture tests were conducted at 649°C and a load stress of 448 MPa. The specimens solution-treated at 982°C were designed with an emphasis on creep rupture strength, and creep rupture tests were conducted at 649°C and a load stress of 483 MPa.

[0025]

[0026]

[0027]

[0028] Example 2: A steel ingot having the composition shown in Table 4 and a steel ingot having the composition shown in Table 5 (both equivalent to SUH660) were prepared as raw materials for hot-worked columnar steel for the present invention. Hot rolling was performed to achieve a total hot working ratio of 630, yielding hot-worked columnar steel with an equivalent circle diameter of 20 mm. Two types of samples were then taken from the surface of the hot-worked steel (before heat treatment) at a depth of D / 2 in the axial direction, one in the axial direction and one perpendicular to the axis, and the GOS values ​​of each sample were determined. The GOS values ​​were calculated in the same manner as in Example 1. Table 6 shows the results of observations of hot-worked steel obtained from a steel ingot having the composition shown in Table 4, and Table 7 shows the results of observations of hot-worked steel obtained from a steel ingot having the composition shown in Table 5. The results in Tables 6 and 7 indicate that the proportion of GOS values ​​of 1.0° or greater in both the longitudinal and transverse sections of the present invention examples was 50% or greater. Next, test specimens (20 mm diameter x 100 mm) were taken in the axial direction from the resulting hot-worked materials, and were subjected to solution treatment and aging treatment to obtain heat-treated materials. Various mechanical properties and grain size numbers were measured. The measurement methods were the same as those used in Example 1. The results of various mechanical properties and grain size numbers are shown in Table 8. Among Samples No. 11-22 and No. 111-118, Samples No. 11, 13, 15, 17, 19, and 21 were obtained from the steel ingots shown in Table 5, while the other samples were obtained from the steel ingots shown in Table 4. The sample solution-treated at 899°C was chosen with emphasis on yield strength, and creep rupture tests of the test specimens were conducted at a temperature of 649°C and a load stress of 448 MPa. The specimens that were solution treated at 982°C were also intended to have improved creep rupture strength, and the creep rupture test conditions for the specimens were a temperature of 649°C and a load stress of 483 MPa.

[0029]

[0030]

[0031]

[0032]

[0033]

[0034] Table 3 for Example 1 and Table 8 for Example 2 show that the samples of the invention exhibited 0.2% proof stress of 590 MPa or more, tensile strength of 1040 MPa, elongation of 22% or more, and reduction of area of ​​40% or more under all processing and heat treatment conditions. The gamma prime phase average equivalent circular diameter was also 10 to 35 nm, demonstrating favorable properties. Furthermore, the samples of the invention also exhibited consistently long and favorable results in creep rupture tests under all processing and heat treatment conditions. On the other hand, the comparative samples exhibited a decrease in 0.2% proof stress (less than 590 MPa) and a shortened time to rupture (less than 30 hours) depending on the processing or heat treatment conditions. The gamma prime phase average equivalent diameter was either greater than 40 nm or less than 9 nm, demonstrating a tendency to be less likely to consistently achieve better mechanical properties and creep rupture properties than the examples of the invention.

[0035] Example 3: A steel ingot (equivalent to SUH660) having the composition shown in Table 9 was prepared as a raw material for hot-worked columnar shapes for the present invention. Hot rolling was performed to achieve a total hot working ratio of 10.1, resulting in a hot-worked columnar material having an area-equivalent circle diameter of 155 mm in a cross section perpendicular to the axis of the columnar material. Two types of samples were then taken from the surface of the hot-worked material (before heat treatment) at positions corresponding to depths D / 4 and D / 2 in the axial direction, one in the axial direction and the other in the transverse direction, and the GOS values ​​of each sample were confirmed. The GOS values ​​were calculated using the same method as in Example 1. Table 10 shows the observation results. From the results in Table 10, the comparative example showed that the proportion of GOS values ​​of 1.0° or more in the longitudinal section was less than 50%, and the difference between the proportion of GOS values ​​of 1.0° or more in the longitudinal section and the transverse section was significant. Next, test specimens (20.5 mm diameter x 100 mm) were taken axially from the surface of the resulting hot-worked material at positions corresponding to depths D / 4 and D / 2 in the axial direction. They were solution-treated and aging-treated to produce heat-treated materials, after which various mechanical properties and grain size numbers were measured. Each measurement method conformed to the corresponding method in Example 1. Table 11 shows the heat treatment conditions for the solution treatment, various mechanical properties, and grain size numbers. Tables 10 and 11 show that, regardless of the heat treatment conditions, Samples No. 121-124, which are comparative examples, all had lower 0.2% yield strengths than the inventive examples of Examples 1 and 2. This is thought to be because the proportion of GOS values ​​of 1.0° or greater in the hot-worked materials of Samples No. 121-124 was less than 50%, meaning that the strain imparted by hot working was insufficient, resulting in the presence of coarse, unrecrystallized grains that were not imparted with processing strain after heat treatment.

[0036]

[0037]

[0038]

Claims

1. A method for producing precipitation hardened austenitic alloy steel, comprising: a hot working step of hot working a material having a precipitation hardened austenitic alloy composition to obtain a hot worked material having an area ratio of 50% or more of crystal grains with a GOS value of 1.0° or more; and a heat treatment step of solution treating the hot worked material and aging treating it to obtain a steel material, wherein the solution treatment temperature in the heat treatment step is 850 to 1050°C, and the aging treatment temperature in the heat treatment step is 700 to 750°C.

2. A precipitation hardened austenitic alloy hot worked material having an area ratio of 50% or more of crystal grains with a GOS value of 1.0° or more, a 0.2% yield strength of 590 MPa or more, a fracture life of 30 hours or more, and a gamma prime phase mean circle equivalent diameter of 10 to 35 nm when solution treated at 850 to 1,050°C and aged at 700 to 750°C.

3. A precipitation hardened austenitic alloy steel having a 0.2% yield strength of 590 MPa or more, a fracture life of 30 hours or more, and an average circular equivalent diameter of gamma prime phase of 10 to 35 nm.

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