Patent Application
20240247331
This application is based on and claims priority under 35 USC 119 from Japanese Patent Application No. 2023-007019 filed on Jan. 20, 2023 and Japanese Patent Application No. 2023-134998 filed on Aug. 22, 2023, the contents of which are incorporated herein by reference.
The present invention relates to austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen and to a manufacturing method therefor. More particularly, the present invention relates to austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen, excellent in hydrogen embrittlement resistance, and relates to a manufacturing method therefor.
In recent years, fuel cell vehicles that use hydrogen as fuel and hydrogen stations that supply hydrogen to fuel cell vehicles have been developed. Various devices used in the fuel cell vehicles, the hydrogen stations, and the like include devices used in a high-pressure hydrogen gas environment (hereinafter, also collectively referred to as “high-pressure hydrogen gas devices”) and devices used in a liquid hydrogen environment (hereinafter, also collectively referred to as “liquid hydrogen devices”). Materials used for such devices are required to have an excellent hydrogen embrittlement resistance. Stainless steel (in particular, austenitic stainless steel with an increased Ni equivalent) has an excellent hydrogen embrittlement resistance and is suitable for this type of application.
Among various kinds of austenitic stainless steel, SUS316L is known as a material excellent in hydrogen embrittlement resistance. Currently, SUS316L is approved as stainless steel excellent in hydrogen embrittlement resistance in accordance with standards of compressed hydrogen container for automobiles stipulated in the Japanese High Pressure Gas Safety Act. However, since SUS316L has low strength, in the case where SUS316L is used in a structural member of a high-pressure hydrogen gas device, the structural member needs to be designed to be thick. As a result, there is a problem that an increase in size and weight of the device cannot be avoided. In order to reduce the weight of the fuel cell vehicle, to make the hydrogen station compact, and to achieve a high-pressure operation in the hydrogen station, it is preferable that the strength and hydrogen embrittlement resistance of the stainless steel used for these applications are high.
In order to solve this problem, various proposals have been made in the related art.
For example, Patent Literature 1 discloses austenitic high-Mn stainless steel containing predetermined amounts of C, N, Si, Cr, Mn, Cu, and Ni and having an austenite stability index Md30 in a predetermined range.
Patent Literature 1 discloses that (A) even in the case of SUS316-based austenitic stainless steel, strain-induced martensite is generated and causes embrittlement in a low-temperature hydrogen environment, and (B) when components are designed such that Md30 satisfies specific conditions, the generation of the strain-induced martensite in a low-temperature hydrogen environment is prevented, and a hydrogen embrittlement resistance sensitivity exceeding that of the SUS316-based austenitic stainless steel is obtained.
Patent Literature 2 discloses austenitic stainless steel containing predetermined amounts of C, Si, Mn, P, S, Ni, Cr, Mo, Cu, N, Al, Ca, O, B, Ti, Nb, and V, in which Creq/Nieq is 1.56 or less, and a P value (an index specifying contents of S, O, and Ca) is −5 or less.
Patent Literature 2 discloses that (A) even in the case of SUS316-based austenitic stainless steel, hydrogen embrittlement may occur under a low-temperature and high-pressure hydrogen gas environment, (B) when the Creq/Nieq is set to 1.56 or less, variations in Ni concentration and reduction in hydrogen embrittlement resistance caused by the variations are prevented, (C) when the amount of S is reduced as much as possible by Al deoxidization and addition of Ca, hot workability can be improved, and (D) accordingly, it is possible to obtain austenitic stainless steel having both excellent hot workability and a hydrogen embrittlement resistance under a low temperature and in a high-pressure hydrogen gas environment of more than 40 MPa.
Patent Literature 3 discloses austenitic stainless steel for high-pressure hydrogen, containing predetermined amounts of C, Si, Mn, P, S, Ni, Cr, and N, in which an area ratio of Cr carbide is 23% or more.
Patent Literature 3 discloses that (A) when the area ratio of the Cr carbide is set to 23% or more, austenitic stainless steel having a 0.2% proof stress of 330 MPa or more and a hardness of 200 Hv or more in a state under a solution heat treatment can be obtained, and (B) when the amount of Ni is set to 8 mass % to 14 mass %, generation of 8 ferrite and generation of deformation-induced martensite can be prevented.
Patent Literature 4 discloses an austenitic stainless steel welded pipe for high-pressure hydrogen transportation, (a) containing predetermined amounts of C, Si, Mn, P, S, Ni, Cr, and N, the balance being Fe and inevitable impurities, (b) having contents of respective elements that are adjusted so as to satisfy a predetermined relational expression, and (c) having an area ratio of a 8 ferrite phase of a welded portion being 0.5% or less.
Patent Literature 4 discloses that hydrogen embrittlement resistance of the austenitic stainless steel welded pipe can be improved by limiting the 8 ferrite phase of the welded portion.
Further, Patent Literature 5 discloses austenitic stainless steel for hydrogen gas, obtained by (a) subjecting austenitic stainless steel containing predetermined amounts of C, Si, Mn, P, S, Cr, Ni, Al, and N and the balance being Fe and impurities to a plastic working at a reduction in cross section of 10% to 50% at room temperature to 200° ° C., and (b) subjecting, in a direction different from the direction of the plastic working, the austenitic stainless steel to another plastic working at a reduction in cross section of 5% or more.
Patent Literature 5 discloses that (A) when the austenitic stainless steel is subjected to a cold working, strength thereof is improved, but when excessive dislocation is introduced into the austenitic stainless steel, a hydrogen embrittlement sensitivity is increased, and (B) when the austenitic stainless steel is subjected to a cold working while changing the working direction, the hydrogen embrittlement sensitivity is significantly reduced.
In a hydrogen station which is a hydrogen supply infrastructure, it is planned to handle hydrogen gas of very low temperature of −40° C. and very high pressure of 82 MPa. In addition, in order to improve efficiency of use of hydrogen gas in the future, a further increase in hydrogen gas pressure is also expected.
However, as disclosed in Patent Literature 1, even SUS316 may become brittle in a low-temperature and high-pressure hydrogen gas environment. In addition, SUS316 is expensive because large amounts of Ni and Mo as rare metals are contained therein.
On the other hand, Patent Literature 2 discloses austenitic stainless steel excellent in hydrogen embrittlement resistance under a hydrogen gas environment of a low temperature and more than 40 MPa. However, Patent Literature 2 does not refer to a high-pressure hydrogen gas environment of about 82 MPa required for a hydrogen station. In addition, in the austenitic stainless steel disclosed in Patent Literature 2, the content of Mo as a rare metal exceeds 2%.
Patent Literature 3 discloses austenitic stainless steel in which deformation-induced martensite transformation at a low temperature is prevented, thereby improving the hydrogen embrittlement resistance in a high-pressure hydrogen gas environment. However, in the austenitic stainless steel disclosed in Patent Literature 3, since the content of Ni as a rare metal is about the same as that in SUS316 in the related art, material cost and resource risk are high.
An object of the present invention is to provide austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen, which has relatively low contents of expensive Ni and Mo, and is excellent in hydrogen embrittlement resistance in a low-temperature and high-pressure hydrogen gas or liquid hydrogen environment.
Another object of the present invention is to provide a manufacturing method for such austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen.
The austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention for solving the above-mentioned objects has the following configurations:
The austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention may further contains, instead of a part of Fe, one or more selected from the group consisting of:
The austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention may further contain, instead of a part of Fe,
A manufacturing method for the austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention includes:
In the case where a relatively large amount of Cu is added to austenitic stainless steel, when heat treatment conditions are inappropriate, Cu is segregated or a low-melting-point Cu compound is easily precipitated. On the other hand, when austenitic stainless steel containing a relatively large amount of Cu is subjected to a soaking treatment at a temperature of 1,200° ° C. or higher, Cu dissolves in an austenite phase, and austenitic stainless steel having less Cu segregation and less low-melting-point Cu compound can be obtained.
Cu is an austenite stabilization element. Therefore, in the case where a relatively large amount of Cu is dissolved in an austenite phase, Nieq becomes large even when a content of Ni or Mo is relatively small, and austenite is stabilized. As a result, generation of a deformation-induced martensite that causes hydrogen embrittlement is prevented, and the hydrogen embrittlement resistance is improved.
Hereinafter, an embodiment of the present invention will be described in detail.
The austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen (hereinafter, also simply referred to as “austenitic stainless steel”) according to the present invention contains the following elements, with the balance being Fe and inevitable impurities. Types of additive elements, component ranges thereof, and reasons for limitation thereof are as follows.
In the present invention, C is an impurity. In the case where the amount of C is excessive, toughness and ductility may be reduced. Therefore, the amount of C needs to be 0.20 mass % or less. The amount of C is preferably 0.1 mass % or less.
The smaller the amount of C, the better. However, extreme reduction in the amount of C causes an increase in manufacturing cost. Considering the manufacturing cost, the amount of C is preferably 0.01 mass % or more.
Si has an effect of improving tensile strength by being solid-dissolved in an austenite phase. In addition, since Si is an austenite stabilization element, Si contributes to an improvement in hydrogen embrittlement resistance. In order to achieve such effects, the amount of Si needs to be 0.10 mass % or more. The amount of Si is preferably 0.50 mass % or more.
On the other hand, in the case where the amount of Si is excessive, grain boundary strength may be reduced, and the hydrogen embrittlement resistance may be reduced. Therefore, the amount of Si needs to be 1.00 mass % or less.
Mn has an effect of forming inclusions such as MnS and improving manufacturability. In addition, since Mn is an austenite stabilization element, Mn contributes to an improvement in hydrogen embrittlement resistance. In order to obtain such effects, the amount of Mn needs to be 0.10 mass % or more. The amount of Mn is preferably 0.50 mass % or more, and more preferably 1.50 mass % or more.
On the other hand, in the case where the amount of Mn is excessive, the solid solubility limit of Cu may be reduced, and a low-melting-point Cu compound may be precipitated. The low-melting-point Cu compound melts during hot working, which causes reduction of hot workability. Therefore, the amount of Mn needs to be 2.0 mass % or less.
In the present invention, P is an impurity. In the case where the amount of P is excessive, hot workability may be reduced. Therefore, the amount of P needs to be 0.050 mass % or less.
The smaller the amount of P, the better. However, extreme reduction in the amount of P causes an increase in manufacturing cost. Considering the manufacturing cost, the amount of P is preferably 0.001 mass % or more.
In the present invention, S is an impurity. In the case where the amount of S is excessive, hot workability may be reduced. Therefore, the amount of S needs to be 0.050 mass % or less. The amount of S is preferably 0.030 mass % or less.
The smaller the amount of S, the better. However, extreme reduction in the amount of S causes an increase in manufacturing cost. Considering the manufacturing cost, the amount of S is preferably 0.001 mass % or more.
Cu has an effect of preventing generation of deformation-induced martensite and improving hydrogen embrittlement resistance. In addition, Cu has an effect of improving cold workability. In order to achieve such effects, the amount of Cu needs to be 2.0 mass % or more. The amount of Cu is preferably 3.0 mass % or more.
On the other hand, in the case where the amount of Cu is excessive, segregation of a component is promoted, and hydrogen embrittlement resistance may be unstable in some portions. In addition, in the case where the amount of Cu is excessive, a Cu concentrated portion may be formed, or Cu exceeding the solid solubility limit may be precipitated as a low-melting-point Cu compound. The Cu concentrated portion or the low-melting-point Cu compound melts during hot working, which causes reduction of hot workability. Therefore, the amount of Cu needs to be less than 4.0 mass %.
Ni has an effect of preventing generation of deformation-induced martensite and improving hydrogen embrittlement resistance. In addition, Ni has an effect of increasing the solid solubility limit of Cu and improving hot workability. In order to achieve such effects, the amount of Ni needs to be 8.0 mass % or more. The amount of Ni is preferably 9.0 mass % or more.
On the other hand, since Ni is expensive, raw material cost increases in the case where the amount of Ni is excessive. Therefore, the amount of Ni needs to be 11.5 mass % or less. The amount of Ni is preferably 10.5 mass % or less.
Cr is a ferrite stabilization element, and in the composition range of the austenitic stainless steel according to the present invention, Cr has an effect of preventing generation of deformation-induced martensite and improving hydrogen embrittlement resistance. In addition, Cr has an effect of enhancing corrosion resistance required for austenitic stainless steel. In order to achieve such effects, the amount of Cr needs to exceed 17.0 mass %. The amount of Cr is preferably 17.5 mass % or more, and more preferably 18.5 mass % or more.
On the other hand, since Cr is expensive, the raw material cost increases in the case where the amount of Cr is excessive. In addition, in the case where the amount of Cr is excessive, formation of 8 ferrite is promoted, and hydrogen embrittlement resistance may be reduced. Therefore, the amount of Cr needs to be 22.0 mass % or less. The amount of Cr is preferably 19.0 mass % or less.
In the present invention, Mo is an element that can be contained as an impurity. Mo is an element that contributes to an improvement of hydrogen embrittlement resistance, but is an expensive element. Therefore, in the case where the amount of Mo is excessive, the raw material cost increases. Therefore, the amount of Mo needs to be 0.20 mass % or less.
In the present invention, the amount of Mo may be zero. However, as described above, Mo contributes to an improvement of hydrogen embrittlement resistance. In order to achieve such an effect, the amount of Mo is preferably 0.01 mass % or more.
In the present invention, N is an impurity. In the case where the amount of N is excessive, stacking fault energy may be reduced, and hydrogen embrittlement resistance may be reduced. Therefore, the amount of N needs to be 0.050 mass % or less.
The smaller the amount of N, the better. However, extreme reduction in the amount of N causes an increase in manufacturing cost. Considering the manufacturing cost, the amount of N is preferably 0.01 mass % or more.
The austenitic stainless steel according to the present invention may contain inevitable impurities.
Here, the “inevitable impurities” refer to components that are mixed in austenitic stainless steel due to various factors such as raw materials and manufacturing processes when the austenitic stainless steel is industrially manufactured, and each of contents thereof are in a range in which the austenitic stainless steel according to the present invention is not adversely affected.
Examples of the inevitable impurities include Sn, Pb, Ti, Ta, and Hf in addition to the above-described C, P, S, Mo, and N. Each of contents of these impurities is preferably 0.05 mass % or less.
Further, the total content of the inevitable impurities is preferably 1.00 mass % or less. The total content is more preferably 0.50 mass % or less.
The austenitic stainless steel according to the present invention may further contain one or two or more of the following elements in addition to the above-described main constituent elements and inevitable impurities. Types of additive elements, component ranges thereof, and reasons for limitation thereof are as follows.
V has an effect of improving strength by generation of a carbonitride. Therefore, the austenitic stainless steel according to the present invention may contain V.
On the other hand, in the case where the amount of V is excessive, a large amount of V-based compound may be generated, and manufacturability may be reduced. Therefore, the amount of V is preferably 0.5 mass % or less.
Nb has an effect of refining crystal grains and improving strength. Therefore, the austenitic stainless steel according to the present invention may contain Nb.
On the other hand, in the case where the amount of Nb is excessive, a large amount of Nb-based compound may be generated, and hot workability may be reduced. Therefore, the amount of Nb is preferably 0.5 mass % or less.
Ca has an effect of improving hot workability and improving manufacturability. Therefore, the austenitic stainless steel according to the present invention may contain Ca.
On the other hand, in the case where the amount of Ca is excessive, a large amount of Ca-based compound may be generated, and manufacturability and corrosion resistance may be reduced. Therefore, the amount of Ca is preferably 0.03 mass % or less.
B has an effect of being segregated at grain boundaries and improving hot workability. Therefore, the austenitic stainless steel according to the present invention may contain B.
On the other hand, in the case where the amount of B is excessive, a large amount of B-based compound may be generated, and workability and corrosion resistance may be reduced. Therefore, the amount of B is preferably 0.05 mass % or less.
Zr is a deoxidizing element. Therefore, the austenitic stainless steel according to the present invention may contain Zr.
On the other hand, in the case where the amount of Zr is excessive, a large amount of Zr-based compound may be generated, and manufacturability may be reduced. Therefore, the amount of Zr is preferably 0.5 mass % or less.
W has an effect of improving strength by generation of a carbide. Therefore, the austenitic stainless steel according to the present invention may contain W.
On the other hand, in the case where the amount of W is excessive, the manufacturing cost and the raw material cost may increase. Therefore, the content of W is preferably 2.0 mass % or less.
Al is a deoxidizing element and has an effect of improving manufacturability. Therefore, the austenitic stainless steel according to the present invention may contain Al.
On the other hand, in the case where the amount of Al is excessive, a large amount of Al-based compound or σ phase may be generated, and manufacturability may be reduced. Therefore, the amount of Al is preferably 0.05 mass % or less. The amount of Al is more preferably 0.01 mass % or less.
Mg has an effect of improving hot workability and improving manufacturability. Therefore, the austenitic stainless steel according to the present invention may contain Mg.
On the other hand, in the case where the amount of Mg is excessive, a large amount of Mg-based compound may be generated, and manufacturability may be reduced. Therefore, the amount of Mg is preferably 0.01 mass % or less.
Co has an effect of preventing generation of deformation-induced martensite and improving hydrogen embrittlement resistance. In the case where Co is contained even in a small amount, such an effect can be obtained. Therefore, the amount of Co is preferably more than 0 mass %. The amount of Co is more preferably 0.01 mass % or more.
On the other hand, in the case where the amount of Co is excessive, ductility is slightly reduced. In addition, Co is an expensive element and the raw material cost is increased. Therefore, the amount of Co is preferably 1.0 mass % or less. The amount of Co is more preferably 0.5 mass % or less, further preferably 0.4 mass % or less, and still more preferably 0.3 mass % or less.
In the present invention, the “Ni equivalent Nieq” refers to a value represented by the following formula (1):
here,
Nieq is an index indicating austenite stability. The larger the Nieq is, the more the deformation-induced martensite transformation in a low-temperature and high-pressure hydrogen gas environment is prevented, and the better the hydrogen embrittlement resistance is. In order to obtain an excellent hydrogen embrittlement resistance, Nieq needs to be 24.0 or more. Nieq is preferably 25.0 or more.
The degree of hydrogen embrittlement resistance can be evaluated based on a magnitude of a relative reduction in area.
Here, the “relative reduction in area” refers to a value represented by the following formula (2). The larger the value of the relative reduction in area represented by the formula (2), the better the hydrogen embrittlement resistance.
here,
In each measurement of A and B, a round bar tensile test piece having a parallel portion diameter of 4 mm is used, and a strain rate thereof is 7×10−5/s.
The austenitic stainless steel according to the present invention has an optimized component, and thus is excellent in hydrogen embrittlement resistance and exhibits high ductility in a low-temperature and high-pressure hydrogen gas environment. In the austenitic stainless steel according to the present invention, in the case where the composition and structure are optimized, the relative reduction in area can be 0.8 or more. In the case where the component and/or structure are further optimized, the relative reduction in area can be 0.9 or more. Such austenitic stainless steel can be obtained by subjecting an ingot having a predetermined composition to a soaking treatment and then to a primary hot working. If necessary, a secondary hot working, a solution treatment, a cold working, or the like may be further performed thereon.
The “tensile strength” refers to a tensile strength obtained by performing a tensile test using a No. 14A test piece having a parallel portion diameter of 6 mm in accordance with JIS Z 2241:2011.
The austenitic stainless steel according to the present invention achieves a tensile strength of 500 MPa or more measured at 25° ° C. by optimizing hot working conditions, solution treatment conditions, and/or cold working conditions. In the case where the component is further optimized, the tensile strength can be 600 MPa or more or 650 MPa or more. In order to obtain austenitic stainless steel having high strength of 650 MPa or more, a cold working is preferably performed.
The “amount of martensite (vol %)” is a value calculated by a method (so-called 5-peak method) using peak intensities of (200) and (211) of a ferrite phase and peak intensities of (200), (220), and (311) of an austenite phase obtained by an X-ray diffraction measurement using a Mo tube, and refers to a value obtained by subtracting a volume fraction (vol %) of the austenite phase calculated based on an integrated intensity ratio of diffraction peaks from 100%.
When the austenitic stainless steel is subjected to a cold working, deformation-induced martensite may be generated. The deformation-induced martensite may reduce hydrogen embrittlement resistance. On the other hand, since the austenitic stainless steel according to the present invention has high austenite stability, the amount of martensite in the metal structure is small even after a cold working.
In the austenitic stainless steel according to the present invention, in the case where the hot working conditions, the solution treatment conditions, and/or the cold working conditions are optimized, the amount of martensite in the metal structure can be 1.0 vol % or less. In the case where manufacturing conditions are further optimized, the amount of martensite can be 0.5 vol % or less.
The austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention may be in any of an as hot-worked state, an as solution-treated state, an as cold-worked state, or a state where a necessary post-processing has been performed after the solution treatment.
In addition, the shape of the austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention is not particularly limited, and an optimum shape can be selected according to the purpose thereof.
Examples thereof include a pipe, a rod, a wire, a plate, and the like.
The austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention is excellent in hydrogen embrittlement resistance, and thus can be used not only for various members exposed to a high-pressure hydrogen gas environment but also for various members exposed to a liquid hydrogen environment.
In particular, since the austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention is excellent in toughness at an extremely low temperature in addition to hydrogen embrittlement resistance, the austenitic stainless steel can be used as a material of a member used in a liquid hydrogen environment, such as (a) a member for a liquid hydrogen pump-pressurized hydrogen station, and (b) a member used in a valve and pump member for liquid hydrogen.
The austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention can be obtained by
First, a raw material blended so as to obtain the austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention is melted and cast to obtain an ingot (first step). Methods and conditions for melting and casting the raw material are not particularly limited, and optimum methods and conditions can be selected according to the purpose thereof. For manufacturing a molten steel, for example, an electric furnace, an argon oxygen decarburization (AOD) furnace, a vacuum oxygen decarburization (VOD) furnace, or the like can be used.
Next, the obtained ingot is subjected to a soaking treatment at a temperature of 1,200° ° C. or higher (second step). The soaking treatment is performed for a purpose of inducing diffusion of components in the steel ingot and removing component segregation.
The temperature of the soaking treatment affects the component segregation. In the case where the soaking treatment is not performed, or in the case where the temperature of the soaking treatment is too low, a Cu concentrated portion may be generated at an interface between a matrix and an oxide scale on a surface, or a low-melting-point Cu compound may be precipitated at a crystal grain boundary. The Cu concentrated portion or the low-melting-point Cu compound causes local melting during a hot working and significantly deteriorates hot workability. In addition, the solid solution temperature of an alloy carbide also increases depending on concentration of Cu. Therefore, in the case where the soaking treatment is not performed, or in the case where the temperature of the soaking treatment is too low, a coarse alloy carbide may remain in a steel material. The coarse alloy carbide causes reduction in toughness and ductility.
In contrast, in the case where the soaking treatment is performed at a high temperature, hot workability and toughness and ductility are improved. In order to achieve such an effect, the temperature of the soaking treatment needs to be 1,200° ° C. or higher.
As the holding time at the temperature of the soaking treatment, an optimum time can be selected according to the purpose thereof. In general, as the holding time at the temperature of the soaking treatment becomes longer, the segregation of Cu is reduced. The optimum holding time varies depending on the temperature of the soaking treatment, but is usually 1 minute to 24 hours. After the holding time is ended, the material is cooled by water cooling, oil cooling, air cooling, or a method capable of achieving a cooling rate equivalent thereto.
In the case where a primary hot working is performed immediately after the soaking treatment, cooling after the soaking treatment can be omitted.
Next, a primary hot working is performed on the material subjected to the soaking treatment (third step). The primary hot working is performed to destroy a coarse cast structure and refine the structure, and at the same time, convert the ingot into steel materials such as slabs, blooms, and billets. The method of the primary hot working is not particularly limited, and an optimum method can be selected according to the purpose thereof. Examples of the method of the primary hot working include a hot forging and a hot rolling.
Next, a secondary hot working is performed on the material after the primary hot working, as necessary (fourth step). The secondary hot working is performed to finish the material obtained in the primary hot working process into a product shape (for example, a steel sheet, a steel bar, a wire material, a steel pipe, or the like) or a shape close thereto. The method of the secondary hot working is not particularly limited, and an optimum method can be selected according to the purpose thereof. Examples of the method of the secondary hot working include a hot rolling, a hot extrusion, and a hot piercing rolling.
Conditions of the secondary hot working are not particularly limited, and optimum conditions can be selected according to the purpose thereof. In addition, the secondary hot working may be performed a plurality of times according to the purpose thereof. The temperature of heating of the steel material performed before the secondary hot working is preferably 900° C. or higher and 1,200° ° C. or lower.
In addition, in the case where the secondary hot working is performed a plurality of times, the temperature of the steel material at completion of the finally performed secondary hot working is preferably 800° ° C. or higher and 1,200° ° C. or lower. This is to optimize crystal grains.
Next, a solution treatment may be performed on the material after the secondary hot working, as necessary (fifth step). The solution treatment may be performed only once, or may be performed a plurality of times.
The temperature of the solution treatment may affect properties of the material. In the case where the solution treatment is not performed, or in the case where the temperature of the solution treatment is too low, a coarse alloy carbide may excessively remain and toughness and ductility may be reduced. Therefore, the temperature of the solution treatment is preferably 1,000° C. or higher.
On the other hand, in the case where the temperature of the solution treatment is too high, there is a possibility that the crystal grains become excessively coarse and strength is reduced. Therefore, the temperature of the solution treatment is preferably 1,200° ° C. or lower, and more preferably 1,150° C. or lower.
As the holding time at the temperature of the solution treatment, an optimum time can be selected according to the purpose thereof. In general, as the holding time of the solution treatment becomes longer, the number density of the coarse alloy carbide becomes smaller. On the other hand, in the case where the holding time is lengthened more than necessary, the crystal grains are excessively coarsened. The optimum holding time varies depending on the temperature of the solution treatment, but is usually 1 minute to 3 hours. After the holding time is ended, the material is cooled by water cooling, oil cooling, air cooling, or a method capable of achieving a cooling rate equivalent thereto.
Next, a cold working may be performed on the material after the secondary hot working or the solution treatment, as necessary (sixth step). The method for the cold working is not particularly limited, and an optimum method can be selected according to the purpose thereof. For example, in the case where a material is subjected to a cold working to obtain a steel pipe, it is preferable to use a cold drawing method. Alternatively, in the case where the material is processed into a steel sheet, it is preferable to use a cold rolling method.
The cold working may affect properties of the material. For the cold working, an optimum working ratio can be selected according to the purpose thereof. In general, as the working ratio of the cold working increases, strength of the material increases. On the other hand, excessive cold working may cause generation of deformation-induced martensite and reduction in hydrogen embrittlement resistance. Therefore, the working ratio of the cold working is preferably 40% or less in terms of reduction in area.
Next, a post-processing may be performed on the material after the primary hot working, the secondary hot working, the solution treatment, or the cold working, as necessary (seventh step). Examples of the post-processing include cutting, welding, and cold working. A member thus obtained is used for various applications.
In the case where a relatively large amount of Cu is added to austenitic stainless steel, when heat treatment conditions are inappropriate, Cu is segregated or a low-melting-point Cu compound is easily precipitated. On the other hand, when austenitic stainless steel containing a relatively large amount of Cu is subjected to a soaking treatment at a temperature of 1,200° ° C. or higher, Cu dissolves in an austenite phase, and austenitic stainless steel having less Cu segregation and less low-melting-point Cu compound can be obtained.
Cu is an austenite stabilization element. Therefore, in the case where a relatively large amount of Cu is dissolved in an austenite phase, Nieq becomes large even when a content of Ni or Mo is relatively small, and austenite is stabilized. As a result, generation of a deformation-induced martensite that causes hydrogen embrittlement is prevented, and the hydrogen embrittlement resistance is improved.
In addition, the austenitic stainless steel in which a relatively large amount of Cu is dissolved is not only excellent in hydrogen embrittlement resistance at a low temperature, but also exhibits excellent workability by which a cold working is possible at a high working ratio and an excellent free-cutting property.
Further, in the austenitic stainless steel according to the present invention, since the stability of austenite is high, deformation-induced martensite is hardly generated. Therefore, in the case where the austenitic stainless steel according to the present invention is subjected to a cold working at a predetermined cold working ratio, high strength can be obtained without reducing hydrogen embrittlement resistance at a low temperature.
In a vacuum induction furnace, 50 kg of steel having a composition shown in Table 1 was melted and cast into an ingot. Thereafter, the ingot (except for Comparative Example 6) was subjected to a soaking treatment at 1,200° ° C. Next, the material after the soaking treatment was subjected to a hot forging. Next, the material after the hot forging was subjected to a solution treatment at 1,080° C. Further, in Examples 2 to 3 and Comparative Example 7, the material after the solution treatment was subjected to a cold working. The reductions in area thereof were 25%, 38%, and 75% in this order, respectively. A test piece was cut out from the obtained material and subjected to various tests. However, in Comparative Example 6, cracking occurred during the hot forging, and the test piece could not be prepared.
The amount of martensite was measured by using a 5-peak method.
A tensile test was performed in accordance with JIS Z 2241:2021.
That is, a round bar tensile test piece (No. 14A test piece) was cut out from the obtained material. The parallel portion of the round bar tensile test piece was parallel to the rolling direction of the steel bar. The diameter of the parallel portion was 6 mm. The tensile strength TS (MPa) was obtained by performing a tensile test on the round bar tensile test piece at normal temperature (25° C.) in the atmosphere.
In order to evaluate hydrogen compatibility, a slow strain rate test was performed. The test temperature was −60° ° C., and the test atmosphere was helium gas or hydrogen gas of 87.5 MPa. A round bar tensile test piece having a parallel portion diameter of 4 mm was used as the test piece. The strain rate was 7×10−5/s.
A reduction in area at rupture in the hydrogen gas and a reduction in area at rupture in the helium gas were separately calculated based on the area of the fracture surface of the round bar tensile test piece after the slow strain rate test. Further, a relative reduction in area at −60° C. (=A/B) was calculated by using them.
Results are shown in Table 2. The following can be found in Table 2.
Regarding the amount of martensite, “A” indicates that the amount of martensite was 0.5 vol % or less, “B” indicates that the amount of martensite was more than 0.5 vol % and 1.0 vol % or less, and “C” indicates that the amount of martensite was more than 1.0 vol %.
Regarding the tensile strength, “A” indicates that the tensile strength was 650 MPa or more, “B” indicates that the tensile strength was 500 MPa or more and less than 650 MPa, and “C” indicates that the tensile strength was less than 500 MPa.
Further, regarding the hydrogen embrittlement resistance, “A” indicates that the relative reduction in area (RRA) was 0.9 or more, “B” indicates that the RRA was 0.8 or more and less than 0.9, and “C” indicates that the RRA was less than 0.8.
(1) In Comparative Example 1, the RRA was less than 0.8. The reason for the decrease in RRA is considered to be that since the amount of Cu was small, generation of deformation-induced martensite in the slow strain rate test caused hydrogen embrittlement.
(2) In Comparative Example 2, the tensile strength after the solution treatment was less than 500 MPa, and the RRA was less than 0.8. The reason for the decrease in tensile strength is considered to be that the amount of solute elements that were dissolved was relatively small (Nieq was less than 24) and sufficient solid solution strengthening did not occur. In addition, the reason for the decrease in RRA is considered to be that since the Nieq was less than 24.0, generation of deformation-induced martensite caused hydrogen embrittlement in the slow strain rate test.
(3) In Comparative Example 3, the RRA was less than 0.8. This is considered to be because the amount of N was excessive.
(4) In Comparative Example 4, the RRA was less than 0.8. This is considered to be because the amount of C was excessive.
(5) In Comparative Example 5, the RRA was less than 0.8. This is considered to be because the amount of Cr was small and the amount of Cu was slightly small.
(6) In Comparative Example 6, cracking occurred during the hot forging, and the test piece could not be prepared. This is considered to be because the soaking treatment was not performed and thus Cu was segregated.
(7) In Comparative Example 7, the amount of martensite exceeded 1.0 vol %, and the RRA was reduced. This is considered to be because the cold working ratio was excessively large, and thus deformation-induced martensite was generated.
(8) In all of Examples 1, 4 to 11, and 13 to 14, the amount of martensite was 0.5 vol % or less, the tensile strength after the solution treatment was 500 MPa or more, and the RRA was 0.9 or more.
(9) In Examples 2 to 3, the tensile strength after the cold working was 650 MPa or more. In addition, in Examples 2 to 3, the amount of martensite was 0.5 vol % or less even after the cold working.
(10) In Example 12, the RRA was slightly reduced. This is considered to be because the amount of Mn and the amount of Cu were slightly small.
In a vacuum induction furnace. 50 kg of steel having a composition shown in Table 3 was melted and cast into an ingot. Thereafter, the ingot was subjected to a soaking treatment at 1.200° ° C. Next, the material after the soaking treatment was subjected to a hot forging. Next, the material after the hot forging was subjected to a solution treatment at 1,080° ° C. Further, in Examples 16 and 17, the material after the solution treatment was subjected to a cold working. The reductions in area thereof were 25% and 38% in this order, respectively. A test piece was cut out from the obtained material and subjected to various tests.
The amount of martensite, the tensile strength, and the hydrogen embrittlement resistance were evaluated in the same manner as in Example 1.
Results are shown in Table 4. The following can be found in Table 4. In Table 4, meanings of “A”, “B”, and “C” of each evaluation item are the same as those in Table 2.
(1) In Comparative Example 8, the RRA was less than 0.8. The reason for the decrease in RRA is considered to be that since the amount of Cu was small, generation of deformation-induced martensite in the slow strain rate test caused hydrogen embrittlement.
(2) In Comparative Example 9, the tensile strength after the solution treatment was less than 500 MPa, and the RRA was less than 0.8. The reason for the decrease in tensile strength is considered to be that the amount of solute elements that were dissolved was relatively small (Nieq was less than 24) and sufficient solid solution strengthening did not occur. In addition, the reason for the decrease in RRA is considered to be that since the Nieq was less than 24.0, generation of deformation-induced martensite caused hydrogen embrittlement in the slow strain rate test.
(3) In Comparative Example 10, the RRA was less than 0.8. This is considered to be because the amount of N was excessive.
(4) In Comparative Example 11, the RRA was less than 0.8. This is considered to be because the amount of C was excessive.
(5) In Comparative Example 12, the RRA was less than 0.8. This is considered to be because the amount of Cr was small and the amount of Cu was slightly small.
(6) In all of Examples 15, 18 to 26, and 28 to 30, the amount of martensite was 0.5 vol % or less, the tensile strength after the solution treatment was 500 MPa or more, and the RRA was 0.9 or more.
(7) In Examples 16 to 17, the tensile strength after the cold working was 650 MPa or more. In addition, in Examples 16 to 17, the amount of martensite was 0.5 vol % or less even after the cold working.
(8) In Example 27, the RRA was slightly reduced. This is considered to be because the amount of Mn and the amount of Cu were slightly small.
Although the embodiments of the present invention have been described in detail above the present invention is not limited to the above embodiments, and various modifications can be made without departing from the gist of the present invention.
The austenitic stainless steel for high-pressure hydrogen gas or liquid hydrogen according to the present invention can be used as a structural member used in a high-pressure hydrogen gas device or a liquid hydrogen device.
| Number | Date | Country | Kind |
|---|---|---|---|
| 2023-007019 | Jan 2023 | JP | national |
| 2023-134998 | Aug 2023 | JP | national |
