EP4756045A1 - Method for producing precipitation hardening austenitic alloy steel material, hot-worked precipitation hardening austenitic alloy material, and precipitation hardening austenitic alloy steel material - Google Patents

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

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Publication number
EP4756045A1
EP4756045A1 EP24849286.0A EP24849286A EP4756045A1 EP 4756045 A1 EP4756045 A1 EP 4756045A1 EP 24849286 A EP24849286 A EP 24849286A EP 4756045 A1 EP4756045 A1 EP 4756045A1
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EP
European Patent Office
Prior art keywords
hot
precipitation hardening
austenitic alloy
heat treatment
steel material
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Pending
Application number
EP24849286.0A
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German (de)
French (fr)
Inventor
Masaru Fujiyoshi
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Proterial Ltd
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Proterial Ltd
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Application filed by Proterial Ltd filed Critical Proterial Ltd
Publication of EP4756045A1 publication Critical patent/EP4756045A1/en
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel with boron
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of annealing
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/004—Heat treatment of ferrous alloys containing Cr and Ni
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/02—Hardening by precipitation
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D7/00—Modifying the physical properties of iron or steel by deformation
    • C21D7/13—Modifying the physical properties of iron or steel by deformation by hot working
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/004—Dispersions; Precipitations

Definitions

  • the present invention relates to a method for producing precipitation hardening austenitic alloy steel material, a precipitation hardening austenitic alloy hot-worked material, and a precipitation hardening austenitic alloy steel material.
  • Precipitation hardening austenitic alloys such as SUH660 are known as members suitable for applications such as bolts and shafts because they have good strength properties over a wide temperature range. Further, since they also excel in hydrogen embrittlement resistance in hydrogen environments, they are known as members suitable for hydrogen station applications.
  • Patent Document 1 discloses that A286 alloy (SUH660 equivalent material), which is a ⁇ ' phase precipitation strengthening alloy, is suitable as an alloy having high mechanical properties relating to hydrogen embrittlement resistance in hydrogen or when hydrogen is absorbed.
  • Patent Document 2 also describes that a forged product is obtained by forging A286 alloy (SUH660 equivalent material) suitable for hydrogen energy equipment at a total forging ratio of 5:1.
  • Non-Patent Document 1 describes that when hydrogen charging was applied to precipitation-strengthened A286 test pieces, properties such as tensile strength of the test pieces decreased with an increase in hydrogen content. From such a background, further strengthening of materials is also required for precipitation hardening austenitic alloys.
  • Non-Patent Document 2 describes that the composition of A286 alloy is improved to achieve higher strength in order to suppress precipitation of ⁇ phase, which increases hydrogen embrittlement susceptibility.
  • Non-Patent Document 2 The method described in Non-Patent Document 2 is effective for improving mechanical properties of high-temperature strength, but there is concern about a decrease in ductility such as elongation accompanying the improvement in strength. Further, in Non-Patent Document 2, W is contained in a relatively large amount for characteristic improvement, but since W is one of the expensive raw materials, it is one of the raw materials whose use is desirable to be avoided as much as possible in industrial application products. Further, in order to accommodate applications for jet engine turbine rotors, heat-resistant bolts, and heat-resistant components for automobiles, improvement in creep rupture strength, which represents deformation resistance in high-temperature environments, is also required in addition to further strengthening.
  • an object of the present invention is to provide a method for producing a precipitation hardening austenitic alloy steel material that may further improve mechanical properties without performing composition improvement of the steel material, and that may be expected to extend the life of components in harsh high-temperature high-load environments and high-pressure hydrogen environments so as to be usable in aircraft, energy, and automobile applications.
  • one aspect of the present invention is a method for producing a precipitation hardening austenitic alloy steel material, the method includes: a hot working step of performing hot working on a material for hot working having a composition of a precipitation hardening austenitic alloy to obtain a hot-worked material having crystal grains with a GOS value of 1.0° or more at an area ratio of 50% or more; and a heat treatment step of performing a solution heat treatment and an aging treatment on the hot-worked material to obtain a steel material, where a solution heat treatment temperature in the heat treatment step is 850 to 1050°C, and an aging treatment temperature in the heat treatment step is 700 to 750°C.
  • Another aspect of the present invention is a precipitation hardening austenitic alloy hot-worked material having crystal grains with a GOS value of 1.0° or more at an area ratio of 50% or more, where when a solution heat treatment at 850 to 1050°C and an aging treatment at a temperature of 700 to 750 are performed, a 0.2% proof stress is 590 MPa or more, a rupture life is 30 hours or more, and a gamma prime phase average equivalent circle diameter is 10 to 35 nm.
  • Another aspect of the present invention is a precipitation hardening austenitic alloy steel material having a 0.2% proof stress of 590 MPa or more, a rupture life of 30 hours or more, and a gamma prime phase average equivalent circle diameter of 10 to 35 nm.
  • the present invention is directed to a precipitation hardening austenitic alloy steel material.
  • This precipitation hardening austenitic alloy steel material refers to SUH309, SUH310, SUH330, SUH660, SUH661 described in JIS-G-4311, and improved materials thereof, and is preferably SUH309, SUH310, SUH330, SUH660, or SUH661. More preferably, it is applied to SUH660.
  • the material contains, in mass%, Ni: 10 to 40% and Cr: 10 to 30%, and that Fe+Ni+Cr is 95% or more in mass%.
  • the material may contain Ni: 10 to 40% and Cr: 10 to 30%, with the balance being Fe and unavoidable impurities.
  • a more preferable lower limit of the Ni content is 20% in mass%, and a more preferable upper limit of the Ni content is 30% in mass%. Further, a more preferable upper limit of Cr is 20% in mass%. Furthermore, in order to improve hardness, high-temperature strength, and the like, one or more elements selected from V, Si, Mn, Al, B, Ta, W, Ti, Mo, and Nb may be contained up to a total of 5.0% in mass%. In addition, examples of unavoidable impurity elements include C, S, P, and O, and it is preferable that the upper limit of each of these is, for example, 0.1%.
  • the material contains Ni: 10 to 40% and Cr: 10 to 30%, contains one or more elements selected from V, Si, Mn, Al, B, Ta, W, Ti, Mo, and Nb up to a total of 5.0% in mass%, with the balance being balance Fe and unavoidable impurities.
  • a material for hot working having a precipitation hardening austenitic alloy composition is prepared.
  • This material for hot working is preferably a steel ingot that may be obtained by casting.
  • hot working in the present embodiment refers to performing hot plastic working such as hot pressing or hot rolling on a steel ingot, and then machining the same into a shape of a round bar or a square bar.
  • this steel billet may be used as a material for forging.
  • remelting may be performed for the purpose of reducing component segregation and non-metallic inclusions.
  • the hot-worked material of the present invention is characterized by having crystal grains with a GOS (Grain Orientation Spread) value of 1.0° or more at an area ratio of 50% or more.
  • GOS Garin Orientation Spread
  • a preferable area ratio is 60% or more, and a more preferable area ratio is 70% or more.
  • This GOS value may be measured by the conventionally known "SEM-EBSD method (electron backscatter diffraction method)" and may be derived by calculating the misorientation of points (pixels) constituting the crystal grains.
  • the crystal misorientation obtained from the GOS value is an index indicating the strain imparted to the alloy by working, and in the case of having crystal grains with a GOS value of 1.0° or more at an area ratio of 50% or more, there is a tendency to easily obtain a uniform fine grain structure with little variation in grain size after solution heat treatment. Further, there are few coarse unrecrystallized grains to which working strain has not been imparted, and there is a tendency to improve mechanical properties, particularly 0.2% proof stress. On the other hand, in the case where 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, and there is a concern that mechanical properties such as 0.2% proof stress may be reduced.
  • the GOS value of the present embodiment may be derived by collecting a sample at a position of D/4 or D/2 in depth in the axial center direction from the surface of the hot-worked material (D is the distance between surfaces passing through the axial center; for example, corresponding to the diameter in the case of a cylindrical shape, or the diagonal length in the case of a prismatic shape) and measuring the same.
  • the cross section for observing the area ratio includes the cross section perpendicular to the axial direction and the cross section in the axial direction, and it is preferable that the area ratio of crystal grains with a GOS value of 1.0° or more is 50% or more in both cases of observation in the cross section perpendicular to the axial direction and the cross section in the axial direction of the bar material.
  • the area ratio of crystal grains with a GOS value of 1.0° or more after the solution heat treatment and the aging heat treatment tends to be a low value of less than 50% because the strain from the hot working is released.
  • the GOS value described above in the present embodiment may be measured by collecting a test piece from the surface of the obtained hot-worked material in the direction as shown in FIG. 1 , at a position of depth D/4 or from depth D/2 in the axial center direction.
  • the position of depth D/2 corresponds to the center of the steel material and is the position where strain is most difficult to be introduced, so it is preferable that the test piece collected from this position of depth D/2 also satisfies the properties specified in the present invention.
  • hot working such that the total hot working ratio during hot working is 15 or more. This allows working strain to be introduced to the center of the hot-worked material, and recrystallization accompanying dynamic recrystallization occurs, thereby improving the mechanical properties of the hot-worked material.
  • Representative examples of "hot working" in the present embodiment include hot forging and hot rolling, and hot forging and hot rolling may also be combined.
  • the total hot working ratio of the present invention indicates the total forging forming ratio (hereinafter also referred to as total forging ratio).
  • total forging forming ratio (total forging ratio)" in the present invention is derived using the forging forming ratio of solid forging described in JIS-G-0701. For example, in the case where hot forging shown in FIG. 13 is performed, the formula ((S 0 /S 1 ) ⁇ (S 1 /S 2 ) ⁇ ⁇ (S n-1 /S n )) becomes the total forging forming ratio.
  • S 0 , S 1 , S 2 ⁇ S n-1 , S n are cross-sectional areas.
  • the total forging forming ratio is 17.3, which satisfies the requirements of the present invention.
  • the upper limit of the total forging ratio is not particularly limited, but since manufacturing costs increase as the number of forging operations increases, for example, 200 is realistic, and 150 may also be used.
  • the cross-sectional area before rolling and the cross-sectional area after rolling are measured and applied to the formula for deriving the forging forming ratio of solid forging described above, which may be used as the hot working ratio.
  • the working ratio may be increased at lower cost compared to hot forging, so the total hot working ratio when hot rolling is applied may be made larger compared to that when hot forging is applied (for example, the total hot working ratio is 1000, etc.).
  • hot forging is selected for hot working unless otherwise specified.
  • the total forging ratio is performed without including the forging forming ratio (upset ratio) during upset forging. For example, in the case where hot forging shown in FIG.
  • 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.
  • Hot forging in the present invention preferably includes hot free forging in which a workpiece is placed on an anvil having a flat or curved surface and the material is processed by pressing or striking with an anvil or die similarly having a flat or curved surface, from the viewpoint of the degree of freedom in the shape to be processed.
  • die forging using a die may be performed, radial forging in which a forged material is obtained by pressing from four directions over the entire length of the material while rotating the material in the circumferential direction may be performed, or processing combining these may be performed.
  • the number of upset forging is preferably at least two times or more in order to introduce working strain to the inside even in large forged products with an area equivalent circle diameter of 100 mm or more.
  • the solution heat treatment and the aging treatment are performed on the hot-worked material obtained by the hot working step to obtain a precipitation hardening austenitic alloy steel material.
  • the solution heat treatment in the present embodiment is set to 850 to 1050°C.
  • the solution heat treatment temperature may be adjusted in a temperature range according to desired properties, and in the case where yield strength is emphasized, the solution heat treatment temperature is preferably 850 to 950°C, and more preferably 870°C to 930°C. In the case where creep rupture strength is emphasized, the solution heat treatment temperature is preferably 950 to 1050°C, and more preferably 970 to 1030°C.
  • the aging treatment temperature in the present embodiment is 700 to 750°C.
  • tensile properties, particularly yield strength may be improved as much as possible, and the creep rupture strength may also be moderately increased.
  • the preferable lower limit of the aging treatment temperature is 710°C, and the more preferable lower limit of the aging treatment temperature is 715°C.
  • the preferable upper limit of the aging treatment temperature is 745°C.
  • the precipitation hardening austenitic alloy steel material of the present invention may be applied to a steel material having a columnar shape and an area equivalent circle diameter of 10 mm or more in a cross section perpendicular to the axial direction.
  • the steel material of the present invention obtained by the producing method of the present invention described above may be applied to materials of various sizes with an area equivalent circle diameter of 10 mm or more.
  • the area equivalent circle diameter is 25 mm or more. More preferably, it is 50 mm or more, 100 mm or more, 200 mm or more, 300 mm or more, and 400 mm or more.
  • annular shape refers to, for example, a cylindrical shape or a prismatic shape, and a taper may be formed in the axial direction.
  • the area equivalent circle diameter of the cross section perpendicular to the axial direction of the present invention may be measured at the cross section perpendicular to the axis at a position of L/2, where L is the length of the steel material.
  • the precipitation hardening austenitic alloy steel material of the present invention obtained by the producing method of the present invention has a 0.2% proof stress of 590 MPa or more.
  • the precipitation hardening austenitic alloy steel material obtained by the producing method of the present invention has a 0.2% proof stress of 590 to 630 MPa.
  • the preferable lower limit of the 0.2% proof stress is 610 MPa.
  • the precipitation hardening austenitic alloy steel material obtained by the producing method of the present invention may have a 0.2% proof stress of 655 MPa or more (preferably 670 MPa or more).
  • elongation is preferably 20% or more, more preferably 23% or more, and further preferably 25% or more.
  • tensile strength is preferably 900 MPa or more, and more preferably 950 MPa or more.
  • reduction of area is preferably 20% or more, and more preferably 30% or more.
  • the precipitation hardening austenitic alloy steel material of the present invention obtained by the producing 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 crystal grains, the hot working ratio also has to be set excessively, and in the case where hot forging is mainly applied, production becomes difficult, so the upper limit may be set to, for example, 10.0.
  • the preferable upper limit of the grain size number is 8.0, and the more preferable upper limit of the grain size number is 7.0.
  • the 0.2% proof stress is affected by the crystal grain size and tends to decrease due to the presence of coarse unrecrystallized grains to which sufficient working strain has not been imparted, but with the steel material of the present invention obtained by the producing method of the present invention, it is possible to achieve a good value without decreasing the 0.2% proof stress.
  • the 0.2% proof stress, elongation, tensile strength, and grain size described above in the present embodiment are measured by collecting a test piece from a position at a depth of D/4 or a position of D/2 in the axial center direction from the surface of the obtained forged material in the direction as shown in FIG. 2 , similarly to the GOS value.
  • the precipitation hardening austenitic alloy steel material of the present invention has high yield strength as described above, while also exhibiting a good value of creep rupture strength.
  • This creep rupture strength indicates the rupture life (rupture time) obtained by performing a creep rupture test defined in ASTM E139.
  • the precipitation hardening austenitic alloy steel material according to the present invention has an excellent rupture life, such as a rupture life of 30 hours or more in the creep rupture test.
  • This rupture life indicates the rupture life in the case where the creep rupture test is performed under conditions of a test temperature of 649°C and a load stress of 448 MPa, and under conditions of a test temperature of 649°C and a load stress of 483 MPa.
  • the precipitation hardening austenitic alloy steel material obtained by the producing method of the present invention maintains high yield strength as described above, while also exhibiting good characteristics in rupture life.
  • the rupture life is 50 hours or more, and preferably, the rupture life is 80 hours or more.
  • the upper limit may be set to 300 hours, for example.
  • the preferable upper limit is 290 hours.
  • the more preferable upper limit is 280 hours.
  • the precipitation hardening austenitic alloy steel material of the present invention has an equivalent circle diameter of the gamma prime phase of 10 nm to 35 nm.
  • the gamma prime phase average equivalent circle diameter is preferably 11.0 nm or more, and preferably 15.0 nm or more.
  • the gamma prime phase average equivalent circle diameter is preferably 34.0 nm or less. In the case where the gamma prime phase equivalent circle diameter is less than 10 nm, depending on the working ratio, there is a tendency that the 0.2% proof stress may not be sufficiently ensured. This is presumed to be because it does not sufficiently contribute to hindering dislocation glide motion inside the steel.
  • the rupture time may significantly decrease. This is presumed to be because dislocation glide motion becomes easier accompanying a change in the mechanism of dislocation glide motion inside the steel.
  • a steel ingot (SUH660 equivalent material) having the composition shown in Table 1 was prepared as a columnar material for hot working for the invention example, and solid forging and upset forging were performed multiple times so that the total hot working ratio became 15.8, and a hot-worked material having an area equivalent circle diameter of 160 mm in the cross section perpendicular to the axis of the columnar material was obtained. Thereafter, two types of samples in the axial direction and the direction perpendicular to the axial direction were collected from the surface of the hot-worked material (before heat treatment) at a depth D/4 position and a D/2 position in the axial center direction, and the GOS value was confirmed for each sample.
  • the GOS value was measured using a field emission scanning electron microscope manufactured by ZEISS and an EBSD measurement and analysis system OIM (Orientation-Imaging-Micrograph) manufactured by TSL, observing the longitudinal cross section (axial direction cross section) and the transverse cross section (cross section perpendicular to the axial direction) of the sample, and measurement samples (11 ⁇ 10 ⁇ 5 mm) were collected from each cross section as shown in FIG. 1 .
  • the measurement surface was 11 ⁇ 10 mm
  • the measurement field of view was 1500 ⁇ m ⁇ 1500 ⁇ m
  • the step distance between adjacent pixels was 3.0 ⁇ m.
  • test pieces (diameter 15 mm ⁇ 100 mm) were collected in the axial direction from a position at a depth of D/4 (D: diameter of area equivalent circle diameter) in the axial center direction, in the direction shown as "material collected from axial direction" in FIG. 2 , and after the solution heat treatment and the aging treatment were applied, various mechanical properties and grain size numbers were measured.
  • the test pieces used for measuring 0.2% proof stress, tensile strength, elongation, and reduction of area were circular cross-section test pieces specified in ASTM E8, and the tests were conducted based on the ASTM E8 tensile test method for metallic materials.
  • the test pieces used for measuring the rupture time were circular cross-section test pieces specified in ASTM E139, and the tests were conducted based on the ASTM E139 creep rupture test method for metallic materials.
  • the grain size number was determined using the standard grain size chart Plate I in accordance with JIS-G-0551.
  • the gamma prime phase average equivalent circle diameter was derived by observing at 100,000 times or 200,000 times magnification using a scanning electron microscope (SEM) and extracting any 10 gamma prime phases from within the field of view. The results of various mechanical properties and grain size numbers are shown in Table 3.
  • the sample with a solution heat treatment temperature of 899°C was one emphasizing yield strength, and the creep rupture test condition for the test piece was conducted at a temperature of 649°C and a load stress of 448 MPa. Further, the sample with a solution heat treatment temperature of 982°C was one also aiming at improvement of creep rupture strength, and the creep rupture test condition for the test piece was conducted at a temperature of 649°C and a load stress of 483 MPa.
  • Table 6 shows the observation results of the hot-worked material obtained from the steel ingot having the composition of Table 4
  • Table 7 shows the observation results of the hot-worked material obtained from the steel ingot having the composition of Table 5. From the results of Table 6 and Table 7, the invention examples have an occupancy rate of 50% or more for GOS values of 1.0° or more in both the longitudinal cross section and the transverse cross section.
  • test pieces (diameter 20 mm ⁇ 100 mm) were collected in the axial direction from the obtained hot-worked material, and the solution heat treatment and the aging treatment were applied to the obtain heat-treated material, after which various mechanical properties and grain size number were measured.
  • Each measurement method followed the method used in Example 1.
  • the results of various mechanical properties and grain size numbers are shown in Table 8.
  • samples No. 11 to 22 and No. 111 to 118 samples No. 11, 13, 15, 17, 19, and 21 are samples obtained from the steel ingot shown in Table 5, and the other data are samples obtained from the steel ingot shown in Table 4.
  • the sample with a solution heat treatment temperature of 899°C was one emphasizing yield strength, and the creep rupture test condition for the test piece was conducted at a temperature of 649°C and a load stress of 448 MPa. Further, the sample with a solution heat treatment temperature of 982°C was one also aiming at improvement of creep rupture strength, and the creep rupture test condition for the test piece was conducted at a temperature of 649°C and a load stress of 483 MPa.
  • Example 3 of Example 1 and Table 8 of Example 2 the samples of the invention examples showed 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 working conditions and heat treatment conditions, and the gamma prime phase average equivalent circle diameter was also 10 to 35 nm in all cases, confirming that they had favorable properties. Further, regarding the rupture life obtained in the creep rupture test, the samples of the invention examples were stably long under all working conditions and heat treatment conditions, showing favorable results.
  • a steel ingot (SUH660 equivalent material) having the composition shown in Table 9 was prepared as a columnar material for hot working for the invention example, and hot rolling was performed so that the total hot working ratio was 10.1, to obtain a hot-worked material having an area equivalent circle diameter of 155 mm in the cross section perpendicular to the axis of the columnar material. Thereafter, two types of samples in the axial direction and the direction perpendicular to the axial direction were collected from positions corresponding to a depth of D/4 and a position corresponding to D/2 in the axial center direction from the surface of the hot-worked material (before heat treatment), and the GOS value was confirmed for each sample.
  • the calculation method of the GOS value is the same as the method used in Example 1.
  • Table 10 shows the observation result. From the results of Table 10, in the comparative example, the occupancy rate of GOS value of 1.0° or more in the longitudinal cross section was less than 50%, and the difference in the occupancy rate of GOS value of 1.0° or more between the longitudinal cross section and the transverse cross section was large.
  • test pieces (diameter 20.5 mm ⁇ 100 mm) were collected in the axial direction from a position corresponding to a depth of D/4 in the axial center direction and a position corresponding to D/2 from the surface of the obtained hot-worked material, and after the solution heat treatment and the aging treatment were applied to obtain heat-treated material, various mechanical properties and grain size number were measured. Each measurement method followed the method used in Example 1. Table 11 shows the heat treatment conditions for the solution heat treatment, various mechanical properties, and grain size numbers.
  • sample Nos. 121 to 124 which are comparative examples, all showed lower 0.2% proof stress compared to the invention examples of Example 1 and Example 2, regardless of the heat treatment conditions. This is considered to be because the occupancy rate of GOS value of 1.0° or more in the hot-worked material of sample Nos. 121 to 124 was less than 50%, and the strain imparted by hot working was insufficient, resulting in the presence of coarse unrecrystallized grains to which working strain was not imparted after heat treatment.

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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

    Technical Field
  • The present invention relates to a method for producing precipitation hardening austenitic alloy steel material, a precipitation hardening austenitic alloy hot-worked material, and a precipitation hardening austenitic alloy steel material.
    • Related Art
  • Precipitation hardening austenitic alloys such as SUH660 are known as members suitable for applications such as bolts and shafts because they have good strength properties over a wide temperature range. Further, since they also excel in hydrogen embrittlement resistance in hydrogen environments, they are known as members suitable for hydrogen station applications. For example, Patent Document 1 discloses that A286 alloy (SUH660 equivalent material), which is a γ' phase precipitation strengthening alloy, is suitable as an alloy having high mechanical properties relating to hydrogen embrittlement resistance in hydrogen or when hydrogen is absorbed. Patent Document 2 also describes that a forged product is obtained by forging A286 alloy (SUH660 equivalent material) suitable for hydrogen energy equipment at a total forging ratio of 5:1.
  • As described in the patent documents mentioned above, precipitation hardening austenitic alloys excel in hydrogen embrittlement resistance, but in high-temperature environments or in high-pressure hydrogen gas, localization of deformation accompanying plastic deformation occurs due to the presence of hydrogen, and stacking faults are easily formed. Such defects tend to cause crack initiation, and mechanical properties such as tensile strength tend to decrease compared to those in air. For example, Non-Patent Document 1 describes that when hydrogen charging was applied to precipitation-strengthened A286 test pieces, properties such as tensile strength of the test pieces decreased with an increase in hydrogen content. From such a background, further strengthening of materials is also required for precipitation hardening austenitic alloys. Non-Patent Document 2 describes that the composition of A286 alloy is improved to achieve higher strength in order to suppress precipitation of η phase, which increases hydrogen embrittlement susceptibility.
    • Citation List Patent Literature
    • Non-Patent Document
    • SUMMARY OF INVENTION Technical Problem
  • The method described in Non-Patent Document 2 is effective for improving mechanical properties of high-temperature strength, but there is concern about a decrease in ductility such as elongation accompanying the improvement in strength. Further, in Non-Patent Document 2, W is contained in a relatively large amount for characteristic improvement, but since W is one of the expensive raw materials, it is one of the raw materials whose use is desirable to be avoided as much as possible in industrial application products. Further, in order to accommodate applications for jet engine turbine rotors, heat-resistant bolts, and heat-resistant components for automobiles, improvement in creep rupture strength, which represents deformation resistance in high-temperature environments, is also required in addition to further strengthening.
    Thus, an object of the present invention is to provide a method for producing a precipitation hardening austenitic alloy steel material that may further improve mechanical properties without performing composition improvement of the steel material, and that may be expected to extend the life of components in harsh high-temperature high-load environments and high-pressure hydrogen environments so as to be usable in aircraft, energy, and automobile applications.
    • Solution to Problem
  • The present invention has been made in view of the above-described problems.
    That is, one aspect of the present invention is a method for producing a precipitation hardening austenitic alloy steel material, the method includes: a hot working step of performing hot working on a material for hot working having a composition of a precipitation hardening austenitic alloy to obtain a hot-worked material having crystal grains with a GOS value of 1.0° or more at an area ratio of 50% or more; and a heat treatment step of performing a solution heat treatment and an aging treatment on the hot-worked material to obtain a steel material, where a solution heat treatment temperature in the heat treatment step is 850 to 1050°C, and an aging treatment temperature in the heat treatment step is 700 to 750°C.
    Another aspect of the present invention is a precipitation hardening austenitic alloy hot-worked material having crystal grains with a GOS value of 1.0° or more at an area ratio of 50% or more, where when a solution heat treatment at 850 to 1050°C and an aging treatment at a temperature of 700 to 750 are performed, a 0.2% proof stress is 590 MPa or more, a rupture life is 30 hours or more, and a gamma prime phase average equivalent circle diameter is 10 to 35 nm.
    Another aspect of the present invention is a precipitation hardening austenitic alloy steel material having a 0.2% proof stress of 590 MPa or more, a rupture life of 30 hours or more, and a gamma prime phase average equivalent circle diameter of 10 to 35 nm.
    • Effects of Invention
  • According to the present invention, it is possible to produce a precipitation hardening austenitic alloy steel material that may further improve mechanical properties without performing composition improvement of the steel material, and that may be expected to extend the life of components in high-temperature high-load environments or high-pressure hydrogen environments.
    • BRIEF DESCRIPTION OF DRAWINGS
    • [FIG. 1] is a schematic diagram showing EBSD sample collection positions in Examples.
    • [FIG. 2] is a schematic diagram showing a sample collection position for a mechanical properties measurement in Examples.
    • [FIG. 3] is a schematic diagram for describing the forging forming ratio.
    • [FIG. 4] is a schematic diagram for describing another forging forming ratio.
    DESCRIPTION OF EMBODIMENTS
  • Hereinafter, the present invention is described in detail. However, the present invention is not limited to the embodiments described herein, and appropriate combinations and improvements are possible without departing from the technical concept of the invention. The present invention is directed to a precipitation hardening austenitic alloy steel material. This precipitation hardening austenitic alloy steel material refers to SUH309, SUH310, SUH330, SUH660, SUH661 described in JIS-G-4311, and improved materials thereof, and is preferably SUH309, SUH310, SUH330, SUH660, or SUH661. More preferably, it is applied to SUH660. Specifically, it is preferable that the material contains, in mass%, Ni: 10 to 40% and Cr: 10 to 30%, and that Fe+Ni+Cr is 95% or more in mass%. Alternatively, the material may contain Ni: 10 to 40% and Cr: 10 to 30%, with the balance being Fe and unavoidable impurities.
  • A more preferable lower limit of the Ni content is 20% in mass%, and a more preferable upper limit of the Ni content is 30% in mass%. Further, a more preferable upper limit of Cr is 20% in mass%. Furthermore, in order to improve hardness, high-temperature strength, and the like, one or more elements selected from V, Si, Mn, Al, B, Ta, W, Ti, Mo, and Nb may be contained up to a total of 5.0% in mass%. In addition, examples of unavoidable impurity elements include C, S, P, and O, and it is preferable that the upper limit of each of these is, for example, 0.1%. Preferably, the material contains Ni: 10 to 40% and Cr: 10 to 30%, contains one or more elements selected from V, Si, Mn, Al, B, Ta, W, Ti, Mo, and Nb up to a total of 5.0% in mass%, with the balance being balance Fe and unavoidable impurities.
  • In the producing method of the present embodiment, first, a material for hot working having a precipitation hardening austenitic alloy composition is prepared. This material for hot working is preferably a steel ingot that may be obtained by casting. Here, hot working in the present embodiment refers to performing hot plastic working such as hot pressing or hot rolling on a steel ingot, and then machining the same into a shape of a round bar or a square bar. In addition, in the case where a machined steel billet (billet, bloom, or the like) is subjected to homogenization heat treatment, this steel billet may be used as a material for forging. In addition, when producing a steel ingot, remelting may be performed for the purpose of reducing component segregation and non-metallic inclusions.
  • Subsequently, the prepared material for hot working is heated in a heating furnace, and multiple hot working operations are performed to obtain a hot-worked material. The hot-worked material of the present invention is characterized by having crystal grains with a GOS (Grain Orientation Spread) value of 1.0° or more at an area ratio of 50% or more. A preferable area ratio is 60% or more, and a more preferable area ratio is 70% or more. This GOS value may be measured by the conventionally known "SEM-EBSD method (electron backscatter diffraction method)" and may be derived by calculating the misorientation of points (pixels) constituting the crystal grains. The crystal misorientation obtained from the GOS value is an index indicating the strain imparted to the alloy by working, and in the case of having crystal grains with a GOS value of 1.0° or more at an area ratio of 50% or more, there is a tendency to easily obtain a uniform fine grain structure with little variation in grain size after solution heat treatment. Further, there are few coarse unrecrystallized grains to which working strain has not been imparted, and there is a tendency to improve mechanical properties, particularly 0.2% proof stress. On the other hand, in the case where 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, and there is a concern that mechanical properties such as 0.2% proof stress may be reduced. The GOS value of the present embodiment may be derived by collecting a sample at a position of D/4 or D/2 in depth in the axial center direction from the surface of the hot-worked material (D is the distance between surfaces passing through the axial center; for example, corresponding to the diameter in the case of a cylindrical shape, or the diagonal length in the case of a prismatic shape) and measuring the same. Further, the cross section for observing the area ratio includes the cross section perpendicular to the axial direction and the cross section in the axial direction, and it is preferable that the area ratio of crystal grains with a GOS value of 1.0° or more is 50% or more in both cases of observation in the cross section perpendicular to the axial direction and the cross section in the axial direction of the bar material. Further, the area ratio of crystal grains with a GOS value of 1.0° or more after the solution heat treatment and the aging heat treatment tends to be a low value of less than 50% because the strain from the hot working is released. In addition, the GOS value described above in the present embodiment may be measured by collecting a test piece from the surface of the obtained hot-worked material in the direction as shown in FIG. 1, at a position of depth D/4 or from depth D/2 in the axial center direction. In particular, the position of depth D/2 corresponds to the center of the steel material and is the position where strain is most difficult to be introduced, so it is preferable that the test piece collected from this position of depth D/2 also satisfies the properties specified in the present invention.
  • In order to obtain the GOS value of the present invention described above, it is preferable to perform hot working such that the total hot working ratio during hot working is 15 or more. This allows working strain to be introduced to the center of the hot-worked material, and recrystallization accompanying dynamic recrystallization occurs, thereby improving the mechanical properties of the hot-worked material. Representative examples of "hot working" in the present embodiment include hot forging and hot rolling, and hot forging and hot rolling may also be combined. In the case where hot forging is selected, the total hot working ratio of the present invention indicates the total forging forming ratio (hereinafter also referred to as total forging ratio). The "total forging forming ratio (total forging ratio)" in the present invention is derived using the forging forming ratio of solid forging described in JIS-G-0701. For example, in the case where hot forging shown in FIG. 13 is performed, the formula ((S0/S1)×(S1/S2)×····· ×(Sn-1/Sn)) becomes the total forging forming ratio. In the above formula, S0, S1, S2·····Sn-1, Sn are cross-sectional areas. As an example of working where the total forging ratio exceeds 15, in the case where a round bar with a diameter of 100 mm is used as a forging material, and solid forging is performed from diameter 100 mm to diameter 50 mm, and then solid forging is performed from diameter 50 mm to diameter 24 mm, the total forging forming ratio is 17.3, which satisfies the requirements of the present invention. The upper limit of the total forging ratio is not particularly limited, but since manufacturing costs increase as the number of forging operations increases, for example, 200 is realistic, and 150 may also be used. In the case where hot rolling is selected for hot working, the cross-sectional area before rolling and the cross-sectional area after rolling are measured and applied to the formula for deriving the forging forming ratio of solid forging described above, which may be used as the hot working ratio. Here, in the case of hot rolling, the working ratio may be increased at lower cost compared to hot forging, so the total hot working ratio when hot rolling is applied may be made larger compared to that when hot forging is applied (for example, the total hot working ratio is 1000, etc.). Hereinafter, in the present embodiment, hot forging is selected for hot working unless otherwise specified.
  • In the case where hot forging is applied to hot working in the present embodiment, it is preferable to perform upset forging at least once so as to increase the total forging ratio without extremely reducing the size of the forged material. In the present embodiment, since the working strain imparted by solid forging tends to primarily contribute to the improvement of mechanical properties, when performing upset forging, in the present embodiment, the calculation of the total forging ratio is performed without including the forging forming ratio (upset ratio) during upset forging. For example, in the case where hot forging shown in FIG. 4 is performed (upset forging between S1 and S2), the formula ((S0/S1)×(S2/S3)×·····×(Sn-1/Sn)) becomes the total forging forming ratio. By combining upset forging, sufficient working strain may be introduced even for large forged products with an area 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.
  • "Hot forging" in the present invention preferably includes hot free forging in which a workpiece is placed on an anvil having a flat or curved surface and the material is processed by pressing or striking with an anvil or die similarly having a flat or curved surface, from the viewpoint of the degree of freedom in the shape to be processed. Certainly, in the case of processing into a complex shape, die forging using a die may be performed, radial forging in which a forged material is obtained by pressing from four directions over the entire length of the material while rotating the material in the circumferential direction may be performed, or processing combining these may be performed.
  • Further, in order to achieve the total forging ratio of the present embodiment described above, it is preferable to perform solid forging at least three times or more. This is because the deformation resistance of precipitation hardening austenitic alloy steel material is large, and thus stable processing may be achieved when the processing rate per forging is smaller. Further, the number of upset forging is preferably at least two times or more in order to introduce working strain to the inside even in large forged products with an area equivalent circle diameter of 100 mm or more.
  • In the present embodiment, the solution heat treatment and the aging treatment are performed on the hot-worked material obtained by the hot working step to obtain a precipitation hardening austenitic alloy steel material. First, the solution heat treatment in the present embodiment is set to 850 to 1050°C. Further, the solution heat treatment temperature may be adjusted in a temperature range according to desired properties, and in the case where yield strength is emphasized, the solution heat treatment temperature is preferably 850 to 950°C, and more preferably 870°C to 930°C. In the case where creep rupture strength is emphasized, the solution heat treatment temperature is preferably 950 to 1050°C, and more preferably 970 to 1030°C.
  • The aging treatment temperature in the present embodiment is 700 to 750°C. By performing aging treatment in the narrow temperature range described in the present invention after performing the hot working step and the solution heat treatment described above, tensile properties, particularly yield strength, may be improved as much as possible, and the creep rupture strength may also be moderately increased. The preferable lower limit of the aging treatment temperature is 710°C, and the more preferable lower limit of the aging treatment temperature is 715°C. The preferable upper limit of the aging treatment temperature is 745°C.
  • The precipitation hardening austenitic alloy steel material of the present invention may be applied to a steel material having a columnar shape and an area equivalent circle diameter of 10 mm or more in a cross section perpendicular to the axial direction. The steel material of the present invention obtained by the producing method of the present invention described above may be applied to materials of various sizes with an area equivalent circle diameter of 10 mm or more. Further, since the production method of the invention may introduce working strain to the center of the material even in large steel material, it is preferable that the area equivalent circle diameter is 25 mm or more. More preferably, it is 50 mm or more, 100 mm or more, 200 mm or more, 300 mm or more, and 400 mm or more. In particular, it is preferable to apply the production method of the present invention to a steel material having an area equivalent circle diameter of 100 mm or more, because it is possible to introduce working strain to the center of the material even in large steel material. Here, "columnar shape" refers to, for example, a cylindrical shape or a prismatic shape, and a taper may be formed in the axial direction. In the case where a taper is formed, the area equivalent circle diameter of the cross section perpendicular to the axial direction of the present invention may be measured at the cross section perpendicular to the axis at a position of L/2, where L is the length of the steel material.
  • The precipitation hardening austenitic alloy steel material of the present invention obtained by the producing method of the present invention has a 0.2% proof stress of 590 MPa or more. In the case of improving the creep rupture strength as well, it is preferable that the precipitation hardening austenitic alloy steel material obtained by the producing method of the present invention has a 0.2% proof stress of 590 to 630 MPa. The preferable lower limit of the 0.2% proof stress is 610 MPa. Further, in the case where a minimum creep rupture strength (for example, a rupture strength of 30 hours or more described later) is acceptable, the precipitation hardening austenitic alloy steel material obtained by the producing method of the present invention may have a 0.2% proof stress of 655 MPa or more (preferably 670 MPa or more). Further, elongation is preferably 20% or more, more preferably 23% or more, and further preferably 25% or more. Further, tensile strength is preferably 900 MPa or more, and more preferably 950 MPa or more. Furthermore, reduction of area is preferably 20% or more, and more preferably 30% or more.
  • The precipitation hardening austenitic alloy steel material of the present invention obtained by the producing 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 crystal grains, the hot working ratio also has to be set excessively, and in the case where hot forging is mainly applied, production becomes difficult, so the upper limit may be set to, for example, 10.0. The preferable upper limit of the grain size number is 8.0, and the more preferable upper limit of the grain size number is 7.0. In particular, the 0.2% proof stress is affected by the crystal grain size and tends to decrease due to the presence of coarse unrecrystallized grains to which sufficient working strain has not been imparted, but with the steel material of the present invention obtained by the producing method of the present invention, it is possible to achieve a good value without decreasing the 0.2% proof stress. The 0.2% proof stress, elongation, tensile strength, and grain size described above in the present embodiment are measured by collecting a test piece from a position at a depth of D/4 or a position of D/2 in the axial center direction from the surface of the obtained forged material in the direction as shown in FIG. 2, similarly to the GOS value.
  • The precipitation hardening austenitic alloy steel material of the present invention has high yield strength as described above, while also exhibiting a good value of creep rupture strength. This creep rupture strength indicates the rupture life (rupture time) obtained by performing a creep rupture test defined in ASTM E139. The precipitation hardening austenitic alloy steel material according to the present invention has an excellent rupture life, such as a rupture life of 30 hours or more in the creep rupture test. This rupture life indicates the rupture life in the case where the creep rupture test is performed under conditions of a test temperature of 649°C and a load stress of 448 MPa, and under conditions of a test temperature of 649°C and a load stress of 483 MPa. The precipitation hardening austenitic alloy steel material obtained by the producing method of the present invention maintains high yield strength as described above, while also exhibiting good characteristics in rupture life. Preferably, the rupture life is 50 hours or more, and preferably, the rupture life is 80 hours or more. The longer the rupture life, the better, but in the case where the rupture life becomes extremely long, tensile properties such as yield strength tend to decrease, so the upper limit may be set to 300 hours, for example. The preferable upper limit is 290 hours. The more preferable upper limit is 280 hours.
  • The precipitation hardening austenitic alloy steel material of the present invention has an equivalent circle diameter of the gamma prime phase of 10 nm to 35 nm. Particularly in the present invention, in order to exhibit good mechanical properties, the gamma prime phase average equivalent circle diameter is preferably 11.0 nm or more, and preferably 15.0 nm or more. Further, the gamma prime phase average equivalent circle diameter is preferably 34.0 nm or less. In the case where the gamma prime phase equivalent circle diameter is less than 10 nm, depending on the working ratio, there is a tendency that the 0.2% proof stress may not be sufficiently ensured. This is presumed to be because it does not sufficiently contribute to hindering dislocation glide motion inside the steel. Further, in the case where the gamma prime phase equivalent circle diameter exceeds 35 nm, depending on the heat treatment condition, the rupture time may significantly decrease. This is presumed to be because dislocation glide motion becomes easier accompanying a change in the mechanism of dislocation glide motion inside the steel.
    • Examples
  • The present invention will be described in more detail in the following examples.
    • (Example 1)
  • A steel ingot (SUH660 equivalent material) having the composition shown in Table 1 was prepared as a columnar material for hot working for the invention example, and solid forging and upset forging were performed multiple times so that the total hot working ratio became 15.8, and a hot-worked material having an area equivalent circle diameter of 160 mm in the cross section perpendicular to the axis of the columnar material was obtained. Thereafter, two types of samples in the axial direction and the direction perpendicular to the axial direction were collected from the surface of the hot-worked material (before heat treatment) at a depth D/4 position and a D/2 position in the axial center direction, and the GOS value was confirmed for each sample. The GOS value was measured using a field emission scanning electron microscope manufactured by ZEISS and an EBSD measurement and analysis system OIM (Orientation-Imaging-Micrograph) manufactured by TSL, observing the longitudinal cross section (axial direction cross section) and the transverse cross section (cross section perpendicular to the axial direction) of the sample, and measurement samples (11×10×5 mm) were collected from each cross section as shown in FIG. 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. Further, observation was performed under the condition that boundaries with an misorientation of 5° or more between adjacent pixels were identified as crystal grain boundaries, and from the obtained GOS value map, the area ratio with respect to the entire observation field of view occupied by crystal grains having a GOS value of 1.0° or more was determined. Table 2 shows the observation result. From the results of Table 2, it can be seen that the obtained hot-worked material has an occupancy rate of 50% or more for GOS values of 1.0° or more in both the longitudinal cross section and the transverse cross section.
  • Subsequently, from the surface of the obtained hot-worked material, test pieces (diameter 15 mm × 100 mm) were collected in the axial direction from a position at a depth of D/4 (D: diameter of area equivalent circle diameter) in the axial center direction, in the direction shown as "material collected from axial direction" in FIG. 2, and after the solution heat treatment and the aging treatment were applied, various mechanical properties and grain size numbers were measured. The test pieces used for measuring 0.2% proof stress, tensile strength, elongation, and reduction of area were circular cross-section test pieces specified in ASTM E8, and the tests were conducted based on the ASTM E8 tensile test method for metallic materials. The test pieces used for measuring the rupture time (rupture life) were circular cross-section test pieces specified in ASTM E139, and the tests were conducted based on the ASTM E139 creep rupture test method for metallic materials. The grain size number was determined using the standard grain size chart Plate I in accordance with JIS-G-0551. The gamma prime phase average equivalent circle diameter was derived by observing at 100,000 times or 200,000 times magnification using a scanning electron microscope (SEM) and extracting any 10 gamma prime phases from within the field of view. The results of various mechanical properties and grain size numbers are shown in Table 3. Here, the sample with a solution heat treatment temperature of 899°C was one emphasizing yield strength, and the creep rupture test condition for the test piece was conducted at a temperature of 649°C and a load stress of 448 MPa. Further, the sample with a solution heat treatment temperature of 982°C was one also aiming at improvement of creep rupture strength, and the creep rupture test condition for the test piece was conducted at a temperature of 649°C and a load stress of 483 MPa. [Table 1]
    (mass%)
    C Si Mn Ni Cr Mo V Al Ti B Balance
    0.04 0.07 0.06 25.42 14.93 1.36 0.30 0.26 2.16 0.0044 Fe and unavoidable impurities
    [Table 2]
    Total hot working ratio Area ratio of crystal grains with GOS value of 1.0° or more [%]
    Longitudinal cross section Transverse cross section
    D/2 D/4 D/2 D/4
    15.8 85.1 77.6 91.4 93.8
    [Table 3]
    Sample No. Total hot working ratio Solution heat treatment temperature [°C] Aging treatment temperature [°C] 02% proof stress [MPa] Tensile strength [MPa Elongation [%], Reduction of area [%] Rupture time [h] Grain size number Gamma prime phase average equivalent circle diameter [nm] Remarks
    1 899 750 794 1059 24.4 43.7 134.4 7.5 33.2
    2 740 811 1081 25.4 48.1 156.9 6.0 29.4
    3 730 798 1093 24.6 48.6 248.3 7.5 24.4
    4 720 793 1110 24.6 48.6 258.9 6.5 20.3
    5 710 755 1105 24.3 46.7 306.8 7.5 14.6
    6 700 760 1107 26.4 50.2 325.2 6.5 11.1 Example of present invention
    7 982 750 711 1040 27.5 49.5 112.9 8.0 32.2
    8 740 716 1053 27.5 50.4 151.9 5.0 27.6
    9 730 695 1067 27.5 50.4 206.0 7.5 22.1
    10 15.8 720 638 1064 29.6 53.9 192.8 4.5 19.5
    11 710 603 1053 30.6 53.4 208.3 5.5 14.5
    12 700 592 1049 33.8 55.0 201.4 6.0 12.9
    101 899 780 700 986 25.6 46.8 24.6 6.5 46.7 Comparative example
    102 760 775 1043 24.5 44.4 87.1 6.0 40.3
    103 680 719 1075 29.0 51.3 404.5 6.5 4.6
    104 660 659 1012 32.5 55.4 435.6 6.5 3.6
    105 982 780 615 974 28.8 50.6 19.7 4.0 45.7
    106 760 682 1022 27.8 49.1 66.8 5.0 40.2
    107 680 464 970 36.6 57.5 5.0 4.4
    108 660 433 900 42.1 61.1 5.0 4.0
    • (Example 2)
  • Steel ingots having the composition shown in Table 4 and steel ingots having the composition shown in Table 5 (both being SUH660 equivalent materials) were prepared as columnar materials for hot working for the invention examples, and hot rolling was performed so that the total hot working ratio became 630, to obtain columnar hot-worked materials having an area equivalent circle diameter of 20 mm. Thereafter, two types of samples in the axial direction and the direction perpendicular to the axial direction were collected from the surface of the hot-worked material (before heat treatment) at a depth of D/2 position in the axial center direction, and the GOS value was confirmed for each sample. The calculation method of the GOS value is the same as the method used in Example 1. Table 6 shows the observation results of the hot-worked material obtained from the steel ingot having the composition of Table 4, and Table 7 shows the observation results of the hot-worked material obtained from the steel ingot having the composition of Table 5. From the results of Table 6 and Table 7, the invention examples have an occupancy rate of 50% or more for GOS values of 1.0° or more in both the longitudinal cross section and the transverse cross section.
  • Subsequently, test pieces (diameter 20 mm × 100 mm) were collected in the axial direction from the obtained hot-worked material, and the solution heat treatment and the aging treatment were applied to the obtain heat-treated material, after which various mechanical properties and grain size number were measured. Each measurement method followed the method used in Example 1. The results of various mechanical properties and grain size numbers are shown in Table 8. Among samples No. 11 to 22 and No. 111 to 118, samples No. 11, 13, 15, 17, 19, and 21 are samples obtained from the steel ingot shown in Table 5, and the other data are samples obtained from the steel ingot shown in Table 4. Here, the sample with a solution heat treatment temperature of 899°C was one emphasizing yield strength, and the creep rupture test condition for the test piece was conducted at a temperature of 649°C and a load stress of 448 MPa. Further, the sample with a solution heat treatment temperature of 982°C was one also aiming at improvement of creep rupture strength, and the creep rupture test condition for the test piece was conducted at a temperature of 649°C and a load stress of 483 MPa. [Table 4]
    (mass%)
    C Si Mn Ni Cr Mo V Al Ti B Balance
    0.04 0.05 0.06 25.37 14.87 1.35 0.28 0.31 2.09 0.0040 Fe and unavoidable impurities
    [Table 5]
    (mass%)
    C Si Mn Ni Cr Mo V Al Ti B Balance
    0.04 0.06 0.08 25.47 14.82 1.36 0.35 0.31 2.10 0.0038 Fe and unavoidable impurities
    [Table 6]
    Total hot working ratio Area ratio of crystal grains with GOS value of 1.0° or more [%]
    Longitudinal cross section Transverse cross section
    D/2 D/4 D/2 D/4
    630 84.4 79.1
    [Table 7]
    Total hot working ratio Area ratio of crystal grains with GOS value of 1.0° or more [%]
    Longitudinal cross section Transverse cross section
    D/2 D/4 D/2 D/4
    630 83.3 76.7
    [Table 8]
    Sample No. Total hot working ratio Solution heat treatment temperature [°C] Aging treatment temperature [°C] 02% proof stress [MPa] Tensile strength [MPa] Elongation [%] Reduction of area [%] Rupture time [h] Grain size number Gamma prime phase average equivalent circle diameter [nm] Remarks
    11 899 750 753 1109 25.0 45.4 31.2 8.0 32.7
    12 740 767 1126 24.7 46.8 47.3 8.0 31.2
    13 730 765 1153 25.4 47.9 57.7 8.5 24.4
    14 720 754 1175 24.9 46.4 83.0 8.5 21.4
    15 710 710 1177 25.4 47.2 82.9 8.5 20.3
    16 700 700 1183 25.9 50.1 96.7 9.5 11.9 Example of present invention
    17 982 750 741 1127 24.0 44.3 53.4 6.5 33.5
    18 740 757 1140 23.5 42.4 67.0 8.0 29.3
    19 730 755 1160 24.1 42.8 91.7 7.0 21.5
    20 630 720 726 1159 23.8 45.2 99.1 7.5 21.2
    21 710 644 1154 27.1 52.4 112.9 6.0 15.0
    22 700 641 1148 27.2 51.4 103.4 7.0 11.3
    111 899 780 651 1012 25.3 47.3 5.1 9.5 51.8 Comparative example
    112 760 732 1084 24.4 45.9 22.4 9.0 41.4
    113 680 664 1154 27.6 51.1 115.0 9.0 8.4
    114 660 640 1111 29.0 52.3 153.2 9.5 6.7
    115 982 780 645 1037 25.0 46.7 8.4 7.0 49.2
    116 760 718 1099 23.8 45.5 30.4 7.0 41.2
    117 680 609 1111 30.1 56.6 108.0 8.0 6.5
    118 660 565 1048 32.6 57.9 57.0 8.0 5.8
  • From Table 3 of Example 1 and Table 8 of Example 2, the samples of the invention examples showed 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 working conditions and heat treatment conditions, and the gamma prime phase average equivalent circle diameter was also 10 to 35 nm in all cases, confirming that they had favorable properties. Further, regarding the rupture life obtained in the creep rupture test, the samples of the invention examples were stably long under all working conditions and heat treatment conditions, showing favorable results. On the other hand, in the samples of the comparative examples, depending on the working conditions or heat treatment conditions, a decrease in 0.2% proof stress (less than 590 MPa) and shortening of rupture time (less than 30 h) occurred, and the gamma prime phase average equivalent diameter was also more than 40 nm or less than 9 nm, confirming a tendency that favorable mechanical properties and creep rupture properties were difficult to stably obtain compared to the invention examples.
    • (Example 3)
  • A steel ingot (SUH660 equivalent material) having the composition shown in Table 9 was prepared as a columnar material for hot working for the invention example, and hot rolling was performed so that the total hot working ratio was 10.1, to obtain a hot-worked material having an area equivalent circle diameter of 155 mm in the cross section perpendicular to the axis of the columnar material. Thereafter, two types of samples in the axial direction and the direction perpendicular to the axial direction were collected from positions corresponding to a depth of D/4 and a position corresponding to D/2 in the axial center direction from the surface of the hot-worked material (before heat treatment), and the GOS value was confirmed for each sample. The calculation method of the GOS value is the same as the method used in Example 1. Table 10 shows the observation result. From the results of Table 10, in the comparative example, the occupancy rate of GOS value of 1.0° or more in the longitudinal cross section was less than 50%, and the difference in the occupancy rate of GOS value of 1.0° or more between the longitudinal cross section and the transverse cross section was large.
  • Subsequently, test pieces (diameter 20.5 mm × 100 mm) were collected in the axial direction from a position corresponding to a depth of D/4 in the axial center direction and a position corresponding to D/2 from the surface of the obtained hot-worked material, and after the solution heat treatment and the aging treatment were applied to obtain heat-treated material, various mechanical properties and grain size number were measured. Each measurement method followed the method used in Example 1. Table 11 shows the heat treatment conditions for the solution heat treatment, various mechanical properties, and grain size numbers.
  • From Table 10 and Table 11, sample Nos. 121 to 124, which are comparative examples, all showed lower 0.2% proof stress compared to the invention examples of Example 1 and Example 2, regardless of the heat treatment conditions. This is considered to be because the occupancy rate of GOS value of 1.0° or more in the hot-worked material of sample Nos. 121 to 124 was less than 50%, and the strain imparted by hot working was insufficient, resulting in the presence of coarse unrecrystallized grains to which working strain was not imparted after heat treatment. [Table 9]
    (mass%)
    C Si Mn Ni Cr Mo V Al Ti B Balance
    0.04 0.05 0.07 25.49 14.84 1.34 0.29 0.25 2.11 0.0042 Fe and unavoidable impurities
    [Table 10]
    Total hot working ratio Area ratio of crystal grains with GOS value of 1.0° or more [%]
    Longitudinal cross section Transverse cross section
    D/2 D/4 D/2 D/4
    10.1 2.7 23.0 85.1 50.4
    [Table 11]
    Sample No. Solution heat treatment temperature [°C] Aging treatment temperature [°C] 0.2% proof stress [MPa] Tensile strength [MPa] Elongation [%] Reduction of area [%] Grain size number Remarks
    121 900 720 577 1018 28.7 43.3 4.5 Comparative example
    122 920 720 560 1014 30.0 43.6 5.0
    123 960 720 573 1008 30.6 44.9 3.0
    124 980 720 583 1010 29.9 40.8 3.5

Claims (3)

  1. A method for producing a precipitation hardening austenitic alloy steel material, the method comprising:
    a hot working step of performing hot working on a material for hot working having a composition of a precipitation hardening austenitic alloy to obtain a hot-worked material having crystal grains with a GOS value of 1.0° or more at an area ratio of 50% or more; and
    a heat treatment step of performing a solution heat treatment and an aging treatment on the hot-worked material to obtain a steel material,
    wherein a solution heat treatment temperature in the heat treatment step is 850 to 1050°C,
    and
    an aging treatment temperature in the heat treatment step is 700 to 750°C.
  2. A precipitation hardening austenitic alloy hot-worked material having crystal grains with a GOS value of 1.0° or more at an area ratio of 50% or more,
    wherein when a solution heat treatment at 850 to 1050°C and an aging treatment at a temperature of 700 to 750 are performed, a 0.2% proof stress is 590 MPa or more, a rupture life is 30 hours or more, and a gamma prime phase average equivalent circle diameter is 10 to 35 nm.
  3. A precipitation hardening austenitic alloy steel material having a 0.2% proof stress of 590 MPa or more, a rupture life of 30 hours or more, and a gamma prime phase average equivalent circle diameter of 10 to 35 nm.
EP24849286.0A 2023-08-02 2024-08-02 Method for producing precipitation hardening austenitic alloy steel material, hot-worked precipitation hardening austenitic alloy material, and precipitation hardening austenitic alloy steel material Pending EP4756045A1 (en)

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