Steel material having excellent fatigue characteristics in hydrogen, method for producing same, steel pipe, and method for producing same

Abstract

The purpose of the present invention is to provide: a steel material which has excellent fatigue characteristics in a high-pressure hydrogen environment and which is suitable for use in a steel structure in a high-pressure hydrogen environment, such as a line pipe for a natural gas (the natural gas is a gas mainly composed of a hydrocarbon such as methane or ethane) containing 100% hydrogen or hydrogen having a hydrogen partial pressure of 1 MPa or more; and a method for producing the steel material. A steel pipe and a manufacturing method thereof. A steel material having excellent fatigue characteristics in hydrogen, the steel material having a specific component composition and a specific structure, and having a crack propagation rate da / dN of 1.0 * 10 <-6 > m * cycle <-1 > or less when a stress expansion coefficient in hydrogen of 1 MPa or more is 20 MPa square root of m.

Description

Technical Field

[0001] The present invention relates to a steel material having excellent fatigue properties in hydrogen, a method for producing the same, a steel pipe, and a method for producing the same. Background Art

[0002] As an existing energy infrastructure, there are pipelines for natural gas transmission. These steels are required to suppress the occurrence of hydrogen-induced cracking in acidic environments. On the other hand, in recent years, hydrogen has received great attention worldwide as a clean energy for building a decarbonized society. Therefore, in order to transport hydrogen in large quantities, the construction of a hydrogen transmission network in which natural gas partially mixed with hydrogen and hydrogen are pressure-transmitted in natural gas pipelines as an alternative is studied. The transmission pressure during operation of these pipelines is assumed to be a high pressure of 1 to 40 MPa, and the pipelines are placed in a high-pressure hydrogen exposure environment. There is a concern that steel used in such an environment will suffer from "hydrogen embrittlement" in which hydrogen intrudes into the steel and the properties deteriorate. Therefore, it is necessary not only to have the high toughness and acid resistance required for conventional pipelines, but also to have the resistance to hydrogen embrittlement required in a hydrogen environment.

[0003] Steel structures used in high-pressure hydrogen environments have always used austenitic stainless steels such as SUS316L, which are less susceptible to hydrogen embrittlement than low-alloy steels. However, austenitic stainless steels such as SUS316L have high steel costs and low strength, so if they are designed to withstand high hydrogen pressures, the wall thickness becomes thicker and the price of the hydrogen structure itself becomes expensive. Therefore, as a steel material for hydrogen steel structures, there is a strong demand for a low-alloy steel material that is lower in cost and can withstand high-pressure hydrogen environments.

[0004] In response to such a demand, for example, the steel for high-pressure hydrogen environment described in Patent Document 1 is steel used in a high-pressure hydrogen environment. By reducing the diffusible hydrogen concentration ratio by Ca / S to less than 1.5 or greater than 11, embrittlement due to diffusible hydrogen is suppressed.

[0005] Patent Document 2 uses a low alloy high strength steel adjusted to a specific component composition, and has greater drawing and elongation values ​​in a 45 MPa hydrogen atmosphere than JIS G3128SHY685NS within a tensile strength range of 900 to 950 MPa in the atmosphere, and is excellent in resistance to high pressure hydrogen environment embrittlement.

[0006] In addition, the low-alloy high-strength steel described in Patent Document 3 is a Cr-Mo-based high-strength low-alloy steel, which is tempered at a relatively high temperature of 560 to 580°C to adjust the grain size after tempering to a grain size of 8.4 or more and a tensile strength of 900 to 950 MPa in an extremely narrow range, thereby becoming a low-alloy high-strength steel with excellent resistance to high-pressure hydrogen environment embrittlement and excellent elongation and drawing properties even in a 45 MPa hydrogen atmosphere.

[0007] In addition, low alloy steel for high pressure hydrogen environment is proposed in Patent Document 4. The low alloy steel described in Patent Document 4 further increases the Mo content compared with the existing steel by adding V, raising the tempering temperature, and utilizing V-Mo carbides, thereby improving the carbide morphology at the grain boundary and greatly improving the resistance to hydrogen environment embrittlement.

[0008] In addition, a steel for a high-pressure hydrogen storage container having excellent hydrogen resistance is proposed in Patent Document 5. According to the technology described in Patent Document 5, by performing a long-term stress relief annealing after normalizing treatment during steel plate manufacturing, MC carbides (Mo, V) C are dispersed and precipitated finely and densely, and hydrogen resistance such as hydrogen embrittlement resistance of the steel is improved.

[0009] In addition, a steel material for high-pressure hydrogen storage is proposed in Patent Document 6. The steel material described in Patent Document 6 is proposed to be a steel material having a metal structure of bainite main structure with an area fraction of 90% or more, and cementite with an average grain size of 50 nm or less and an average aspect ratio of 3 or less dispersed in the bainite.

[0010] It should be noted that Non-Patent Document 1 describes the value of fatigue strength of low alloy steel.

[0011] Prior art literature

[0012] Patent Literature

[0013] Patent Document 1: Japanese Patent Application Laid-Open No. 2005-2386

[0014] Patent Document 2: Japanese Patent Application Publication No. 2009-46737

[0015] Patent Document 3: Japanese Patent Application Publication No. 2009-275249

[0016] Patent Document 4: Japanese Patent Application Publication No. 2009-74122

[0017] Patent Document 5: Japanese Patent Application Publication No. 2010-37655

[0018] Patent Document 6: Japanese Patent Application Publication No. 2012-107332

[0019] Non-patent literature

[0020] Non-patent document 1: Matsunaga et al., Int J Hydrogen Energy, Vol. 40 (2015), p. 5739-5748 Summary of the invention

[0021] The pressure inside the pipeline pipe repeatedly loads stress on the structure such as the pipeline pipe due to fluctuations during operation and periodic closures. Therefore, fatigue fracture needs to be considered when designing steel structures such as pipeline pipes. However, as shown in Non-Patent Document 1, it is known that the fatigue life of materials is reduced in a high-pressure hydrogen environment. That is, it means that when designing pipeline pipes based on the previous natural gas pipeline pipes, the service life of the pipeline pipes is reduced. However, although the above-mentioned prior art can suppress the occurrence of hydrogen-induced cracking in an acidic environment, it cannot fully improve the fatigue strength in hydrogen, that is, there is the following problem: it is difficult to obtain high fatigue strength in hydrogen that is more likely to affect the service life in addition to suppressing the occurrence of hydrogen-induced cracking in an acidic environment.

[0022] In view of the above-mentioned problems of the prior art, the present invention aims to provide a steel material having excellent fatigue characteristics in a high-pressure hydrogen environment, a manufacturing method thereof, a steel pipe and a manufacturing method thereof, which is suitable for use as a steel structure used in a high-pressure hydrogen environment, such as a pipeline pipe for 100% hydrogen or natural gas containing hydrogen with a hydrogen partial pressure of 1 MPa or more (natural gas is a gas with hydrocarbons such as methane and ethane as main components).

[0023] It should be noted that the so-called "excellent fatigue properties in a high-pressure hydrogen environment" here means that the crack growth rate da / dN is 1.0×10-10 at the stress enlargement coefficient range = 20 MPa√m obtained by fatigue testing in accordance with ASTM E647 with a frequency of 1 Hz, repetitive waveform: sine wave, control method: load control, and stress ratio: R = 0.1 under two environments: room temperature (20 ± 10 ° C) and a hydrogen pressure of 1 MPa or more or a natural gas (mainly composed of hydrocarbons such as methane and ethane) containing hydrogen with a hydrogen partial pressure of 1 MPa or more. －6 m·cycle －1 The natural gas containing hydrogen at a hydrogen partial pressure of 1 MPa or more refers to, for example, a gas having a hydrogen concentration of 30% or less in volume fraction and a total gas pressure of 30 MPa or less.

[0024] It should be noted that if the crack growth rate da / dN in a hydrogen environment is 1.0×10 －6 m·cycle －1 Next, it is possible to design structural steel for hydrogen use within a range of plate thickness that can be manufactured by the manufacturing process.

[0025] From the above viewpoints, the present inventors have conducted intensive studies on the conditions that various steel materials should satisfy in hydrogen, and have discovered a new steel material and steel pipe having excellent fatigue properties in hydrogen.

[0026] The present invention has been made based on the above-mentioned new findings and further studies, and the gist of the present invention is as follows.

[0027] [1] A steel material having excellent fatigue properties in hydrogen, comprising, by mass%, 0.02 to 0.15% C, 0.05 to 0.5% Si, 0.3 to 2.0% Mn, 0.01 to 0.15% Al, 0.0005 to 0.008% N, 0.03% or less P, 0.01% or less S, 0.01% or less O, more than 0% to 2.5% Nb, 0.0010% or less H % or less, Cu: 0-2.5%, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.03%, Sb: 0-0.3%, Ca: 0-0.01%, Mg: 0-0.01%, REM: 0-0.005%, the remainder is composed of Fe and unavoidable impurities;

[0028] Nb precipitates with equivalent circle diameter of 2nm to 100nm are 10 / μm 2 above,

[0029] The crack growth rate da / dN when the stress expansion coefficient in hydrogen above 1MPa = 20MPa√m is 1.0×10 －6 m·cycle -1 the following.

[0030] [2] A method for manufacturing a steel material, comprising:

[0031] The heating step is to heat the slab having the composition described in [1] at 1000 to 1250° C.

[0032] In the hot rolling step, the slab heated in the heating step is rolled at a finishing temperature of Ar3 point or higher.

[0033] a controlled cooling step of cooling the steel plate obtained in the hot rolling step at an average cooling rate of 10°C / s or more at a temperature at the center of the plate thickness of 1000°C to 400°C and a cooling stop temperature of 250°C to 650°C; and

[0034] The dehydrogenation treatment step is to perform a dehydrogenation treatment on the steel sheet obtained in the above-mentioned controlled cooling step.

[0035] [3] A steel pipe having excellent fatigue properties in hydrogen, comprising, by mass%, 0.02 to 0.15% C, 0.05 to 0.5% Si, 0.3 to 2.0% Mn, 0.01 to 0.15% Al, 0.0005 to 0.008% N, 0.03% or less P, 0.01% or less S, 0.01% or less O, more than 0% to 2.5% Nb, 0.0010 % or less, Cu: 0-2.5%, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.03%, Sb: 0-0.3%, Ca: 0-0.01%, Mg: 0-0.01%, REM: 0-0.005%, the remainder is composed of Fe and unavoidable impurities;

[0036] Nb precipitates with equivalent circle diameter of 2nm to 100nm are 10 / μm 2 above,

[0037] The crack growth rate da / dN when the stress expansion coefficient in hydrogen above 1MPa = 20MPa√m is 1.0×10 －6 m·cycle -1 the following.

[0038] [4] A method for manufacturing a steel pipe, comprising:

[0039] The heating step is to heat the slab having the composition described in [3] at 1000 to 1250°C.

[0040] In the hot rolling step, the slab heated in the heating step is rolled at a finishing temperature of Ar3 point or higher.

[0041] A controlled cooling step is performed to cool the hot rolled steel sheet obtained in the hot rolling step at an average cooling rate of 10°C / s or more at a temperature at the center of the sheet thickness of 1000°C to 400°C and a cooling stop temperature of 250°C to 650°C.

[0042] Any of the following pipe making processes: a pipe making process in which, after the controlled cooling process, the hot rolled steel sheet is bent, both ends are butted and welded; and a pipe making process in which, after the controlled cooling process, the hot rolled steel sheet is formed into a cylindrical shape by cold roll forming, both circumferential ends of the cylindrical shape are butted and resistance welded, and

[0043] The dehydrogenation treatment step is to perform a dehydrogenation treatment on the steel pipe obtained in the pipe making step.

[0044] According to the present invention, a steel material and a steel pipe having extremely excellent fatigue properties in hydrogen in a high-pressure hydrogen environment can be obtained, and the present invention is extremely useful industrially. DETAILED DESCRIPTION

[0045] Next, a method for implementing the present invention will be described in detail. It should be noted that the following description shows a preferred embodiment of the present invention, and the present invention is not limited by the following description. As a first embodiment, a steel material is described in detail, and then as a second embodiment, a UOE steel pipe as an example of a steel pipe of the present invention is described in detail, and as a third embodiment, an electric resistance welded steel pipe as an example of a steel pipe of the present invention is described in detail.

[0046] First embodiment

[0047] [Ingredients]

[0048] The reasons for limiting the chemical composition of the steel material of the present invention are described below. It should be noted that "%" in the following description means "mass %" unless otherwise specified.

[0049] C: 0.02~0.15%

[0050] C is contained to ensure moderate hardenability, but its effect is insufficient when it is less than 0.02%. Therefore, the C content is 0.02% or more. The C content is preferably 0.03% or more. On the other hand, if it exceeds 0.15%, the toughness of the base material and the heat-affected zone of the weld deteriorates, and the weldability deteriorates significantly. Therefore, the C content is 0.15% or less. The C content is preferably 0.12% or less. The C content is more preferably 0.10% or less, and further preferably 0.08% or less.

[0051] Si: 0.05-0.5%

[0052] Si is contained as a deoxidizing material in the steelmaking stage and as an element to ensure hardenability. If it is less than 0.05%, its effect is insufficient. Therefore, the Si content is 0.05% or more. The Si content is preferably 0.1% or more. The Si content is more preferably 0.15% or more. On the other hand, if it exceeds 0.5%, the grain boundaries become embrittled, deteriorating the low-temperature toughness and fatigue properties in hydrogen. Therefore, the Si content is 0.5% or less. The Si content is preferably 0.4% or less. The Si content is more preferably 0.3% or less, and further preferably 0.25% or less.

[0053] Mn: 0.3~2.0%

[0054] Mn is contained as an element to ensure hardenability, and its effect is insufficient when it is less than 0.3%. Therefore, it is 0.3% or more. The Mn content is preferably 0.4% or more. The Mn content is more preferably 0.5% or more. The Mn content is further preferably 0.6% or more. On the other hand, if it contains more than 2.0%, the grain boundary strength decreases and the low-temperature toughness deteriorates. In addition, the hardness of the surface layer and the center segregation portion increases during controlled cooling, so the fatigue properties in hydrogen deteriorate. Therefore, the Mn content is 2.0% or less. The Mn content is preferably 1.8% or less. The Mn content is more preferably 1.5% or less, and further preferably 1.3% or less.

[0055] Al: 0.01~0.15%

[0056] Al is contained as a deoxidizing material, and also has the effect of pinning austenite grains and inhibiting grain coarsening during heating as a fine precipitate of Al-based nitrides. When it is less than 0.01%, its effect is insufficient. Therefore, the Al content is 0.01% or more. The Al content is preferably 0.02% or more. The Al content is more preferably 0.03% or more. On the other hand, if it contains more than 0.15%, the cleanliness of the steel decreases, and the toughness and fatigue properties in hydrogen deteriorate. Therefore, the Al content is 0.15% or less. The Al content is preferably 0.12% or less. The Al content is more preferably 0.10% or less, and further preferably 0.08% or less.

[0057] N: 0.0005~0.008%

[0058] N forms fine precipitates by forming nitrides with Nb, Ti, Al, etc., and has the effect of suppressing the coarsening of grains and improving low-temperature toughness by pinning austenite grains during heating, so it is contained. When it contains less than 0.0005%, the microstructure refinement effect cannot be fully achieved. Therefore, the N content is 0.0005% or more. The N content is preferably 0.001% or more. The N content is more preferably 0.0025% or more. On the other hand, when it contains more than 0.008%, the amount of solid-solution N increases, which will damage the toughness of the base material and the heat-affected zone of the weld, and the fatigue properties in hydrogen will deteriorate. Therefore, the N content is 0.008% or less. The N content is preferably 0.007% or less. The N content is more preferably 0.006% or less, and further preferably 0.005% or less.

[0059] P: 0.03% or less

[0060] P as an impurity element is easy to segregate at the grain boundary. If it exceeds 0.03%, the bonding strength of adjacent grains is reduced, and the low temperature toughness and fatigue properties in hydrogen are deteriorated. Therefore, the P content is 0.03% or less. The P content is preferably 0.02% or less, and more preferably 0.01% or less. The lower limit is not particularly limited, but from the perspective of increasing costs, it is preferably 0.001% or more.

[0061] S: 0.01% or less

[0062] S as an impurity element is easy to segregate at the grain boundary, and MnS as a non-metallic inclusion is also easy to generate. If it exceeds 0.01%, the bonding strength of adjacent grains decreases, the amount of inclusions increases, and the low-temperature toughness and fatigue properties in hydrogen deteriorate. Therefore, the S content is 0.01% or less. The S content is preferably 0.008% or less. The S content is more preferably 0.005% or less, and further preferably 0.002% or less. The lower limit is not particularly limited, but from the perspective of increasing costs, it is preferably 0.0001% or more. The S content is more preferably 0.001% or more.

[0063] O: 0.01% or less

[0064] O affects the workability of the material by forming oxides with Al and the like, so the less the better. When it contains more than 0.01%, inclusions increase and workability is impaired. In addition, as inclusions increase, fatigue properties in hydrogen also deteriorate. Therefore, the O content is 0.01% or less. The O content is preferably 0.008% or less, and more preferably 0.005% or less. The lower limit is not particularly limited, but from the perspective of increasing costs, it is preferably 0.0001% or more. The O content is more preferably 0.001% or more.

[0065] Nb: more than 0% and less than 2.5%

[0066] Nb has the effect of improving hardenability, and as a fine precipitate of Nb-based carbonitride, it pins austenite grains during heating and inhibits the coarsening of grains. The smaller the grain size, the greater the grain boundary area, and the better the fatigue properties in hydrogen. Therefore, the Nb content exceeds 0%. The Nb content is preferably 0.005% or more. The Nb content is more preferably 0.01% or more. On the other hand, when the content exceeds 2.5%, the toughness of the weld heat affected zone deteriorates. Therefore, the Nb content is 2.5% or less. The Nb content is more preferably 2.2% or less. The Nb content is further preferably 2.0% or less, and most preferably 1.5% or less.

[0067] H: 0.0010% or less

[0068] H is sometimes introduced into steel in various processes during manufacturing. If the amount introduced is large, the risk of cracking after solidification increases and fatigue crack propagation is accelerated. In addition, when the amount introduced is large, the crack propagation rate increases, so it is important to reduce the amount of hydrogen in the steel. If it is less than 0.0010%, these effects will not be a problem, so the H content is less than 0.0010%. Preferably it is less than 0.0005%. More preferably, it is less than 0.0002%. On the other hand, less than 0.00001% becomes the main reason for the increase in cost, so the H content is preferably more than 0.00001%. The H content is more preferably more than 0.0001%. It should be noted that the hydrogen content is the amount of residual hydrogen after the steel, steel pipe, UOE, etc. are formed. In the present invention, the above-mentioned H content can be achieved by implementing a dehydrogenation treatment process.

[0069] In the present invention, the remainder of the above-mentioned component composition is preferably a steel composition consisting of Fe and inevitable impurities, but depending on the desired characteristics, it is preferably further contained one or more of Cu: 0-2.5%, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.03%, Sb: 0-0.3%, Ca: 0-0.01%, Mg: 0-0.01%, and REM: 0-0.005%, either alone or in combination.

[0070] Cu: 0-2.5%

[0071] Cu has the effect of improving hardenability. Therefore, when Cu is contained, the Cu content can be 0% or more, but when it is less than 0.05%, it is difficult to obtain the above effect, so the Cu content is preferably 0.05% or more. If the Cu content exceeds 2.5%, hot cracks are easily generated when the steel sheet is heated or welded. Therefore, when Cu is contained, it is 2.5% or less. The Cu content is preferably 2.3% or less. The Cu content is more preferably 2.0% or less, and further preferably 1.8% or less.

[0072] Ni: 0-2.5%

[0073] Ni has the same effect of improving hardenability as Cu, and further has the effect of improving toughness. Therefore, when Ni is contained, the Ni content can be more than 0%, but when it is less than 0.05%, it is difficult to obtain the above effect, so the Ni content is preferably more than 0.05%. If it exceeds 2.5%, the economic efficiency is poor. Therefore, when Ni is contained, the Ni content is less than 2.5%. The Ni content is preferably less than 2.3%. The Ni content is more preferably less than 2.0%, preferably less than 1.8%.

[0074] Cr: 0~2.5%

[0075] Cr is an element that ensures hardenability. When Cr is contained, the Cr content may be 0% or more. However, when it is less than 0.1%, it is difficult to obtain the above effect, so the Cr content is preferably 0.1% or more. On the other hand, when it is contained in excess of 2.5%, weldability deteriorates. Therefore, when Cr is contained, the Cr content is 2.5% or less. The Cr content is preferably 2.3% or less. The Cr content is more preferably 2.0% or less, further preferably 1.5% or less, and most preferably 1.2% or less.

[0076] Mo: 0~2.0%

[0077] Mo has the effect of improving hardenability, so when Mo is contained, the Mo content can be 0% or more, but when it is less than 0.05%, it is difficult to obtain the above effect, so the Mo content is preferably 0.05% or more. On the other hand, when it is contained over 2.0%, the economic efficiency is poor. Therefore, when Mo is contained, the Mo content is 2.0% or less. The Mo content is preferably 1.8% or less. The Mo content is more preferably 1.5% or less, and further preferably 1.2% or less.

[0078] V: 0~0.5%

[0079] V has the effect of improving hardenability, and as a fine precipitate of V-based carbide, it pins austenite grains during heating and inhibits the coarsening of grains. Therefore, when V is contained, the V content can be 0% or more, but when it is less than 0.005%, it is difficult to obtain the above effect, so the V content is preferably 0.005% or more. The V content is preferably 0.01% or more. On the other hand, when the added content exceeds 0.5%, the toughness of the weld heat affected zone is deteriorated. Therefore, when V is contained, the V content is 0.5% or less. The V content is preferably 0.4% or less. The V content is more preferably 0.3% or less, and further preferably 0.2% or less.

[0080] Ti: 0~0.5%

[0081] Ti has the effect of improving hardenability, and has the effect of pinning austenite grains and inhibiting grain growth when heated as a fine precipitate of Ti-based carbonitride. Therefore, when Ti is contained, the Ti content can be 0% or more, but if it is less than 0.005%, it is difficult to obtain the above effect, so the Ti content is preferably 0.005% or more. On the other hand, when the added content exceeds 0.5%, the toughness of the weld heat affected zone deteriorates. Therefore, when Ti is contained, the Ti content is 0.5% or less. The Ti content is preferably 0.4% or less. The Ti content is more preferably 0.3% or less, and further preferably 0.2% or less.

[0082] W: 0~2.5%

[0083] W has the effect of improving hardenability. Therefore, when W is contained, the W content may be 0% or more. However, when it is less than 0.05%, it is difficult to obtain the above effect. Therefore, the W content is preferably 0.05% or more. On the other hand, if it exceeds 2.5%, weldability deteriorates. Therefore, when W is contained, the W content is 2.5% or less. The W content is preferably 2.3% or less. The W content is more preferably 2.0% or less, and further preferably 1.8% or less.

[0084] B: 0~0.005%

[0085] B is contained as an element to ensure hardenability, so when B is contained, the B content can be 0% or more, but when it is less than 0.0005%, it is difficult to obtain the above effect, so the B content is preferably 0.0005% or more. On the other hand, if it exceeds 0.005%, the toughness is deteriorated. Therefore, when B is contained, the B content is 0.005% or less. The B content is preferably 0.004% or less. The B content is more preferably 0.003% or less, and further preferably 0.002% or less.

[0086] Sn: 0~0.03%

[0087] Sn has the effect of improving the corrosion resistance of steel. Therefore, when Sn is contained, the Sn content can be 0% or more, but when it is less than 0.005%, it is difficult to obtain the above effect, so the Sn content is preferably 0.005% or more. On the other hand, when the added content exceeds 0.03%, the high temperature ductility decreases, and the possibility of cracking during casting is increased. Therefore, when Sn is contained, the Sn content is 0.03% or less. The Sn content is preferably 0.025% or less. The Sn content is more preferably 0.02% or less, and further preferably 0.015% or less.

[0088] Sb: 0~0.3%

[0089] Sb has the effect of improving the corrosion resistance of steel. Therefore, when Sb is contained, the Sb content can be 0% or more, but when it is less than 0.005%, it is difficult to obtain the above effect, so the Sb content is preferably 0.005% or more. The Sb content is more preferably 0.01% or more. On the other hand, when the added content exceeds 0.3%, the high temperature ductility decreases and the hot rolling property decreases. Therefore, when Sb is contained, the Sb content is 0.3% or less. The Sb content is preferably 0.25% or less. The Sb content is more preferably 0.2% or less, and further preferably 0.15% or less.

[0090] Ca: 0~0.01%

[0091] Ca forms CaS, which has the effect of controlling the morphology of sulfide inclusions to spherical inclusions that are not easily elongated by rolling, that is, CaS, instead of inclusions that are easily elongated by rolling, that is, MnS. Therefore, when Ca is contained, the Ca content can be 0% or more, but when it is less than 0.0005%, it is difficult to obtain the above effect, so the Ca content is preferably 0.0005% or more. The Ca content is preferably 0.001% or more. On the other hand, when the content exceeds 0.01%, the cleanliness decreases, so the material such as toughness deteriorates. Therefore, when Ca is contained, the Ca content is 0.01% or less. The Ca content is preferably 0.005% or less. The Ca content is more preferably 0.003% or less, and further preferably 0.002% or less.

[0092] Mg: 0~0.01%

[0093] Mg is sometimes used as a molten iron desulfurization material. Therefore, when Mg is contained, the Mg content can be 0% or more, but when it is less than 0.0005%, it is difficult to obtain the above effect, so the Mg content is preferably 0.0005% or more. The Mg content is preferably 0.001% or more. On the other hand, when the added content exceeds 0.01%, it leads to a decrease in cleanliness. Therefore, when Mg is contained, the Mg content is 0.01% or less. The Mg content is preferably 0.005% or less. The Mg content is more preferably 0.004% or less, and further preferably 0.003% or less.

[0094] REM: 0~0.005%

[0095] REM reduces the amount of solid-solubilized S at the grain boundaries and improves the SR crack resistance by generating sulfides as REM (O, S) in the steel. Therefore, when REM is contained, the REM content can be 0% or more, but when it is less than 0.0005%, it is difficult to obtain the above effect, so the REM content is preferably 0.0005% or more. On the other hand, when the added content exceeds 0.005%, REM sulfides accumulate significantly in the precipitated crystal bands, resulting in deterioration of the material. Therefore, when REM is contained, the REM content is 0.005% or less. The REM content is preferably 0.003% or less. The REM content is more preferably 0.001% or less. It should be noted that REM is the abbreviation of Rare Earth Metal, which is a rare earth metal.

[0096] In the chemical composition of the steel plate and the steel pipe, the remainder other than the above-mentioned components (elements) is composed of Fe and inevitable impurity elements.

[0097] The metal structure of the present invention will be described in detail.

[0098] The number of Nb precipitates with a circle equivalent diameter of 2nm to 100nm is 10 / μm 2 above

[0099] The pinning effect of Nb can be confirmed by obtaining a predetermined amount of precipitates. In the present invention, the Nb precipitates having a size of 2 nm to 100 nm in terms of circle equivalent diameter are 10 / μm. 2 The number of Nb precipitates with a circle equivalent diameter of 2 nm to 100 nm is preferably 15 / μm. 2 More than 20 pieces / μm 2 More than 25 / μm is more preferred 2 As for the upper limit, the number of Nb precipitates with a circle equivalent diameter of 2 nm to 100 nm is preferably 100 / μm because the hydrogen embrittlement resistance deteriorates when the number of precipitates increases too much. 2 More preferably, 90 pieces / μm 2 Below, more preferably 80 / μm 2 Below. For example, Nb precipitates include NbC, NbN, Nb carbonitrides, etc. As described later, the cooling stop temperature is important for the precipitation of fine Nb precipitates. It should be noted that the reason for focusing on Nb precipitates of 2 nm or more is because it is difficult to confirm, and the reason for focusing on Nb precipitates of 100 nm or less is because if the Nb precipitates become too coarse, it will have an adverse effect on hydrogen embrittlement.

[0100] The structure of the present invention is not particularly limited, but the main structure is bainite, and the bainite is preferably 60% or more, more preferably 80% or more, and even more preferably 90% or more. In addition, the bainite may be 100%.

[0101] The crack growth rate da / dN when the stress expansion coefficient in hydrogen above 1MPa = 20MPa√m is 1.0×10 －6 m·cycle -1 the following

[0102] The fatigue crack growth rate is an important parameter in the design of gas containers such as pipelines and gas containers, and is necessary to obtain a service life that ensures the safety of fracture structural components. In fracture structural components, it is difficult to make cracks and crack generation positions zero. When subjected to repeated stress, cracks will inevitably occur and expand. Therefore, from the perspective of service life, it is also important to control the crack growth rate of the steel used in the above fracture structural components. The crack growth rate is small when the stress state applied to the front end of the crack is small, and it increases as the stress state at the front end of the crack increases. In a hydrogen environment, hydrogen invades the steel, making cracks easy to develop. The degree of promotion of crack growth rate due to hydrogen is greatly affected by the structure and precipitates of the material. In the crack growth test in hydrogen above 1MPa, if the fatigue test is carried out according to ASTM E647 with a frequency of 1Hz, repetitive waveform: sine wave, control method: load control, and stress ratio: R=0.1, the crack growth rate da / dN when the stress expansion coefficient = 20MPa√m is obtained is 1.0×10 －6 m·cycle -1 Below, the service life of the steel structure in a high-pressure hydrogen environment can be ensured. Therefore, in the crack growth test in hydrogen above 1MPa, the crack growth rate da / dN when the stress expansion coefficient = 20MPa√m is 1.0×10 －6 m·cycle -1 The crack growth rate da / dN when the stress expansion coefficient is 20 MPa√m is preferably 0.9×10 －6 m·cycle -1 Below, more preferably 0.8×10 －6 m·cycle -1 Below, more preferably 0.7×10 －6 m·cycle -1 For the lower limit, it can be considered that it is higher than the result in the atmosphere and the closer it is to the result in the atmosphere, the better. Since the crack growth rate da / dN when the stress expansion coefficient in the atmosphere is 20MPa√m is 0.05×10 －6 m·cycle -1 Therefore, the crack growth rate da / dN when the stress expansion coefficient is 20MPa√m only needs to be 0.05×10 －6 m·cycle -1 That’s all.

[0103] In addition, the plate thickness of the steel material is not particularly limited, but the plate thickness is preferably 4 mm or more. The plate thickness is more preferably 5 mm or more. In addition, the plate thickness is preferably 70 mm or less. The plate thickness is more preferably 30 mm or less.

[0104] The steel material having excellent fatigue properties in hydrogen of the present invention is not particularly limited as long as it has the above-mentioned component composition and metal structure and satisfies the crack growth rate in hydrogen.

[0105] Hereinafter, a thick plate as a steel material having excellent fatigue properties in hydrogen of the present invention is exemplified, and a method for producing the steel material having excellent fatigue properties in hydrogen of the present invention is described.

[0106] The steel material of the present invention can be produced by sequentially performing a heating step, a hot rolling step, a controlled cooling step, and a dehydrogenation treatment step.

[0107] It should be noted that the steel with excellent fatigue properties in hydrogen of the present invention is a steel having the above-mentioned component composition, including various classifications such as thin plates, thick plates, steel pipes, etc. with excellent fatigue crack growth resistance in hydrogen, or can be a steel for hydrogen pipeline formed into a specified shape.

[0108] In the following description of the manufacturing method, unless otherwise specified, the temperature is the temperature at the center of the plate thickness of the steel billet, steel material, or steel pipe. The temperature at the center of the plate thickness of the steel material or the like can be obtained by calculating the temperature distribution in the cross section of the steel plate or the like by heat transfer analysis, and correcting the result by the surface temperature of the steel plate or the like. It should be noted that the above-mentioned "hot-rolled steel plate" also includes hot-rolled plates and hot-rolled steel strips.

[0109] In the present invention, the smelting method of the steel billet (steel billet) is not particularly limited. For example, known smelting methods such as converters, electric furnaces, and vacuum melting furnaces are suitable. The casting method is also not particularly limited. For example, a steel billet of a desired size can be manufactured by a known casting method such as a continuous casting method. It should be noted that there is no problem in applying the ingot casting-opening rolling method instead of the continuous casting method. The molten steel may also be further subjected to secondary refining such as ladle refining.

[0110] Heating process

[0111] Heating temperature: 1000℃~1250℃

[0112] When the heating temperature is less than 1000°C, the deformation resistance of the rolled material becomes large and it is difficult to roll. Therefore, the heating temperature is above 1000°C. The preferred heating temperature is above 1050°C, more preferably above 1100°C, and further preferably above 1120°C. On the other hand, if the heating temperature exceeds 1250°C, the austenite grains coarsen, and fine austenite grains cannot be obtained in the subsequent rolling (rough rolling, finish rolling), and the fatigue resistance in hydrogen is reduced. Therefore, the heating temperature in the hot rolling process is below 1250°C. The above heating temperature is preferably below 1230°C. The above heating temperature is more preferably below 1210°C, and further preferably below 1200°C.

[0113] It should be noted that in the present invention, in addition to the existing method of temporarily cooling the steel billet (slab) to room temperature and then heating it again after manufacturing, it is also possible to apply energy-saving direct rolling processes such as loading it into a heating furnace in a warm sheet state without cooling it to room temperature, or rolling it immediately after slightly keeping it warm.

[0114] Hot rolling process

[0115] After the above heating, hot rolling including rough rolling and finish rolling is performed, and the finish rolling is performed under the following conditions. Hot rolling is performed using a hot rolling mill.

[0116] Finish rolling end temperature: above Ar3 point

[0117] When the finishing temperature of the finishing rolling is less than the Ar3 point, the surface temperature of the steel sheet becomes below the ferrite transformation start temperature during the finishing rolling, and a processed ferrite with a high dislocation density is generated, and the fatigue characteristics in hydrogen are reduced. Therefore, the finishing temperature of the finishing rolling is above the Ar3 point. When the Ar3 point is lower than 770°C, the finishing temperature of the finishing rolling is preferably above 770°C. When the Ar3 point is higher than 770°C, the finishing temperature of the finishing rolling is preferably above Ar3 point + 30°C, and more preferably above Ar3 point + 50°C. On the other hand, the upper limit of the finishing temperature of the finishing rolling is not particularly specified, but when the Ar3 point is lower than 850°C, if it exceeds 850°C, sometimes the reduction in the austenite non-recrystallization temperature range is insufficient, and fine austenite grains cannot be obtained, and the fatigue resistance in hydrogen is reduced. Therefore, when the Ar3 point is lower than 850°C, the finishing temperature of the finishing rolling is preferably below 850°C. More preferably, it is below 830°C. When the Ar3 point is higher than 850°C, the finishing temperature of the finishing rolling is preferably below Ar3 point - 30°C, and more preferably below Ar3 point - 50°C.

[0118] In the present invention, the final plate thickness (plate thickness of the steel plate after finish rolling) is preferably 4 mm or more. On the other hand, the upper limit of the final plate thickness is not particularly specified, but is preferably 70 mm or less from the viewpoint of temperature management of the steel plate.

[0119] The Ar3 point varies depending on the alloy composition of the steel, so it can be obtained by experimentally measuring the phase transition temperature of each steel, but it can also be obtained from the following formula based on the component composition.

[0120] Ar3(℃)=910-310C(%)-80Mn(%)-20Cu(%)-15Cr(%)

[0121] -55Ni(%) -80Mo(%)

[0122] The content of each alloy element is (mass %).

[0123] Controlled cooling process

[0124] Average cooling rate at the center of plate thickness at 1000-400°C: 10°C / s or more

[0125] When the average cooling rate in the center of the plate thickness is less than 10°C / s, the nucleation frequency of ferrite and bainite decreases, and they coarsen, so the fatigue properties in hydrogen decrease. Therefore, the average cooling rate in the center of the plate thickness of 1000-400°C is 10°C / s or more. It is preferably 12°C / s or more, more preferably 15°C / s or more, and further preferably 18°C / s or more. On the other hand, the upper limit of the above average cooling rate is not specifically specified, but if it exceeds 60°C / s, a large amount of hard structure is generated on the surface of the steel plate, and the steel structure having the target structure in the present invention cannot be obtained, and the fatigue properties in hydrogen decrease. Therefore, the average cooling rate in the center of the plate thickness is preferably less than 60°C / s. The average cooling rate in the center of the plate thickness is more preferably less than 55°C / s, and further preferably less than 50°C / s.

[0126] Cooling stop temperature at the center of plate thickness: 250℃~650℃

[0127] When the cooling stop temperature at the center of the plate thickness is less than 250°C, the cooling stop temperature at the surface of the steel plate becomes too low, a large amount of hard structure is generated on the surface of the steel plate, and the steel structure having the target structure in the present invention cannot be obtained, and the fatigue properties in hydrogen are reduced. Therefore, the cooling stop temperature at the center of the plate thickness is 250°C or more. Preferably, it is 280°C or more. More preferably, it is 300°C or more. More preferably, it is 390°C or more, and most preferably, it is 480°C or more. On the other hand, if the cooling stop temperature at the center of the plate thickness exceeds 650°C, the nucleation frequency of ferrite or bainite decreases, and they coarsen, so that the structure having the target average crystal grain size in the present invention cannot be obtained, the strength decreases, and the coarsening of Nb precipitates due to residual heat is aggravated. Therefore, the cooling stop temperature at the center of the plate thickness is 650°C or less. The cooling stop temperature at the center of the plate thickness is preferably 620°C or less, more preferably 600°C or less, and more preferably 580°C or less. Most preferably, it is 550°C or less. More preferably, it is 500°C or less.

[0128] It should be noted that when manufacturing a thin steel plate, it is necessary to coil it into a coil after cooling stops. For the same reasons as above, the temperature of coiling into a coil after cooling stops is very important. The coiling temperature is 650°C or less. Preferably it is 620°C or less, more preferably 600°C or less, and further preferably 580°C or less. Most preferably it is 550°C or less. More preferably it is 500°C or less. On the other hand, if the coiling temperature is too low, coiling becomes difficult or a steel structure with a target structure cannot be obtained, so it is 250°C or more. Preferably it is 280°C or more. More preferably it is 300°C or more. More preferably it is 390°C or more, and most preferably it is 480°C or more.

[0129] In addition, the cooling stop temperature of the steel plate surface is not particularly limited, but when it is less than 250°C, a large amount of hard structure is generated on the steel plate surface, and the steel structure with the target structure fraction in the present invention cannot be obtained, and the fatigue resistance in hydrogen is reduced. Therefore, the cooling stop temperature of the steel plate surface is preferably above 250°C. The cooling stop temperature of the steel plate surface is more preferably above 280°C. On the other hand, if the cooling stop temperature exceeds 650°C, the cooling stop temperature in the center of the plate thickness becomes too high, and the nucleation frequency of ferrite or bainite in the center of the plate thickness decreases, and they coarsen, so the structure with the target average crystal grain size in the present invention cannot be obtained, and the strength is reduced. Therefore, the cooling stop temperature of the plate surface is preferably below 650°C. The cooling stop temperature of the steel plate surface is more preferably below 470°C.

[0130] In the present invention, the average cooling rate is a value (cooling rate) obtained by ((plate thickness center temperature of steel plate before cooling - plate thickness center temperature of hot-rolled steel plate after cooling) / cooling time) unless otherwise specified.

[0131] Examples of the cooling method include water cooling such as spraying water from a nozzle, cooling by spraying a cooling gas, etc. In the present invention, it is preferred to perform cooling operation (treatment) on both surfaces of the steel sheet so that both surfaces are cooled under the same conditions.

[0132] Dehydrogenation process

[0133] When hydrogen is already present in the steel, the acceleration of fatigue crack growth increases, and the fatigue life and fatigue crack growth rate in hydrogen decrease. Therefore, in order to release the hydrogen remaining after manufacturing, dehydrogenation treatment is required. Dehydrogenation treatment can reduce the amount of hydrogen in the steel by keeping it at high temperature for a certain period of time before the product is used, and can obtain a steel plate with excellent fatigue properties in hydrogen under a high-pressure hydrogen environment.

[0134] The holding time R (sec) is preferably determined based on the plate thickness and tube thickness t (mm) of the steel material and the steel tube and the hydrogen diffusion coefficient D (mm·sec) in the steel at room temperature. －1 ) and adopt the following formula (A).

[0135] R≥t 2 / D···(A)

[0136] The hydrogen diffusion coefficient also varies depending on the components and metal structure contained, but for example, the hydrogen diffusion coefficient can be 1×10 －5 ~5×10 －3 mm 2 / s. More preferably, 5×10 －4 mm 2 / s or less.

[0137] The dehydrogenation process is implemented before the welding construction of pipe making or connecting steel pipes. It should be noted that since the hydrogen diffusion coefficient D at high temperature becomes smaller, in order to make hydrogen escape quickly, the dehydrogenation treatment is preferably at high temperature. In the case of high temperature, the diffusion coefficient D' (diffusion coefficient at each temperature) of the temperature at which the value of D of the above (A) formula is maintained can be used for calculation. On the other hand, when the temperature T of the dehydrogenation process is too high, the material strength is significantly reduced, so the dehydrogenation temperature is preferably below 550°C. The dehydrogenation temperature T is more preferably below 500°C. The dehydrogenation temperature T is further preferably below 400°C, and most preferably below 300°C. In addition, since the dehydrogenation treatment at a temperature lower than room temperature is the main reason for the increase in processing time and cost, for this reason, the dehydrogenation temperature T is preferably above room temperature. The dehydrogenation temperature T is more preferably above 50°C. The dehydrogenation temperature T is further preferably above 100°C, and most preferably above 150°C. The so-called dehydrogenation temperature T here is the temperature of the atmosphere in the dehydrogenation process. Room temperature refers to 20±10°C.

[0138] In particular, when heating is performed, it takes time for the temperature Tc of the center of the plate thickness of the steel material and the steel pipe to reach the temperature of the atmosphere (dehydrogenation treatment temperature T) in the dehydrogenation treatment process. Therefore, even if the above-mentioned holding time R (sec) is satisfied at the atmosphere temperature, if the dehydrogenation treatment temperature T (atmosphere temperature) is not reached at the center of the plate thickness, the dehydrogenation treatment may not be sufficient. Therefore, it is preferred to maintain the above-mentioned holding time R (sec) or more after the temperature Tc of the center of the plate thickness reaches the target dehydrogenation treatment temperature T. Furthermore, in order to obtain a specified crack growth rate in hydrogen, it is necessary to appropriately adjust the amount of hydrogen in the steel material at the surface and the center of the plate thickness. For this purpose, it is preferred to maintain the R (sec) specified in formula (A) or more at the dehydrogenation treatment temperature T, and it is further preferred to maintain the above-mentioned holding time R (sec) or more after the temperature Tc of the center of the plate thickness reaches the target dehydrogenation treatment temperature T. In other words, at least the former can properly control the amount of hydrogen in the steel material at the surface of the steel material and the steel pipe, and when the latter is implemented, the amount of hydrogen in the steel material from the surface to the center of the plate thickness of the steel material and the steel pipe can be properly controlled. The plate thickness center temperature Tc may be measured using a thermocouple or the like, or may be estimated using a finite element method or the like.

[0139] It should be noted that the time and temperature of the dehydrogenation treatment process may include the temperature and time applied when heating is performed in the pipe making process of electric resistance welded pipe, UOE, etc. as described later. Furthermore, the scale on the surface of the steel hinders dehydrogenation, so it is preferable to remove the scale and perform the dehydrogenation treatment. The removal method is not limited, for example, it can be physical cleaning using high pressure cleaning, or a chemical method using a scale remover can be used. If about 100 μm is removed in thickness, the effect of scale removal can be obtained.

[0140] Second embodiment

[0141] Hereinafter, a UOE steel pipe will be specifically described as a steel pipe of the present invention. The component composition, metal structure, and crack growth rate of the steel pipe are the same as those described in the steel material, and the manufacturing method, heating step, hot rolling step, controlled cooling step, and dehydrogenation treatment step are also implemented in the same manner as those described in the steel material. The pipe manufacturing step is implemented by the following operation.

[0142] Pipe making process

[0143] The UOE steel pipe of the present invention is a hot-rolled steel plate subjected to bending processing, specifically, the end of the hot-rolled steel plate is subjected to groove processing, and after being formed into a steel pipe shape by a C-type press, a U-type press, or an O-type press, the butt joint is seam welded by inner surface welding and outer surface welding, and further manufactured by a pipe expansion process as required. In addition, any welding method can be used as long as it is a method that can obtain sufficient joint strength and joint toughness, but submerged arc welding is preferably used from the viewpoint of excellent welding quality and manufacturing efficiency. In addition, the steel pipe obtained by seam welding the butt joint after being formed into a tubular shape by press bending can also be expanded. The welding method is not particularly limited, and submerged arc welding and the like can be cited.

[0144] In addition, the plate thickness of the UOE steel pipe cited as an example of the steel pipe of the present invention is preferably 10 mm to 50 mm.

[0145] Third embodiment

[0146] Hereinafter, an electric resistance welded steel pipe will be described as an example of the steel pipe of the present invention. The component composition, metal structure, and crack growth rate of the electric resistance welded steel pipe are the same as those described in the steel material, and the manufacturing method, heating step, hot rolling step, and dehydrogenation treatment step are also carried out in the same manner as those described in the steel material. The controlled cooling step and the pipe making step after hot rolling are carried out by the following operations.

[0147] Controlled cooling process

[0148] Average cooling rate at the center of plate thickness at 1000-400°C: 10°C / s or more

[0149] When the average cooling rate at the center of the plate thickness of 1000-400°C is less than 10°C / s, the nucleation frequency of ferrite and bainite decreases, and they coarsen, so the fatigue properties in hydrogen decrease. Therefore, the average cooling rate at the center of the plate thickness of 1000-400°C is 10°C / s or more. It is preferably 12°C / s or more, more preferably 15°C / s or more, and further preferably 18°C / s or more. On the other hand, the upper limit of the above average cooling rate is not specifically specified, but if it exceeds 60°C / s, a large amount of hard structure is generated on the surface of the steel plate, and the steel structure having the target structure in the present invention cannot be obtained, and the fatigue properties in hydrogen decrease. Therefore, the average cooling rate at the center of the plate thickness is preferably 60°C / s or less. The average cooling rate at the center of the plate thickness is more preferably 55°C / s or less, and further preferably 50°C / s or less.

[0150] Cooling stop temperature at the center of plate thickness: 250℃~650℃

[0151] When the cooling stop temperature at the center of the plate thickness is less than 250°C, the cooling stop temperature at the surface of the steel plate becomes too low, a large amount of hard structure is generated on the surface of the steel plate, and the steel structure having the target structure in the present invention cannot be obtained, and the fatigue properties in hydrogen are reduced. Therefore, the cooling stop temperature at the center of the plate thickness is 250°C or more. Preferably, it is 280°C or more. More preferably, it is 300°C or more. The cooling stop temperature at the center of the plate thickness is more preferably 390°C or more. In order to reliably suppress the formation of hard structure on the surface of the steel plate, the cooling stop temperature at the center of the plate thickness is most preferably 450°C or more. More preferably, it is 480°C or more. On the other hand, if the cooling stop temperature exceeds 650°C, the nucleation frequency of ferrite or bainite decreases, and they coarsen, so the structure having the target average crystal grain size in the present invention cannot be obtained, and the strength is reduced. In addition, the coarsening of Nb precipitates is aggravated due to residual heat. Therefore, the cooling stop temperature at the center of the plate thickness is 650°C or less. Preferably, it is 620°C or less, more preferably, it is 600°C or less, and more preferably, it is 580°C or less. The most preferred temperature is 550°C or lower, and the most preferred temperature is 500°C or lower.

[0152] Then, the hot-rolled steel sheet is coiled into a coil. The coiling temperature is preferably 650° C. or less.

[0153] In addition, the cooling stop temperature of the steel plate surface is not particularly limited, but when it is less than 250°C, a large amount of hard structure is generated on the steel plate surface, and the steel structure having the target structure in the present invention cannot be obtained, and the fatigue properties in hydrogen are reduced. Therefore, on the other hand, if the cooling stop temperature of the steel plate surface exceeds 650°C, the cooling stop temperature in the center of the plate thickness becomes too high, and the nucleation frequency of ferrite or bainite in the center of the plate thickness decreases, and they coarsen, so the structure having the target structure in the present invention cannot be obtained, and the strength is reduced. The cooling stop temperature of the steel plate surface is preferably 280°C or more, preferably 470°C or less.

[0154] It should be noted that, in the present invention, the average cooling rate is a value (cooling rate) obtained by ((center temperature of the hot-rolled steel sheet thickness before cooling - center temperature of the hot-rolled steel sheet thickness after cooling) / cooling time) unless otherwise specified.

[0155] Examples of the cooling method include water cooling by spraying water from a nozzle, cooling by spraying a cooling gas, etc. In the present invention, it is preferred that both surfaces of the steel sheet be cooled under the same conditions.

[0156] Pipe making process

[0157] The electric resistance welded steel pipe cited as an example of the steel pipe of the present invention is manufactured by forming the hot rolled steel plate into a cylindrical shape by cold rolling, and butting and welding the circumferential ends of the cylindrical shape (pipe making process). Furthermore, it is also possible to manufacture the steel pipe by forming the steel pipe into an electric resistance welded steel pipe material using a sizing roll satisfying the following formula (1) (sizing process), and applying an internal pressure p (MPa) satisfying the following formula (2) to the inner surface of the electric resistance welded steel pipe material (internal pressure loading process).

[0158] It should be noted that the above-mentioned cylindrical shape means that the tube circumferential cross section is a "C" shape.

[0159] Diameter of sizing roller (mm) ≥ thickness of hot-rolled steel plate (mm) / 0.020···(1)

[0160] X<p≤X×1.5···(2)

[0161] It should be noted that X = (wall thickness of electric resistance welded steel tube billet (mm) / radius of electric resistance welded steel tube billet (mm)) × yield strength of electric resistance welded steel tube billet (MPa)

[0162] The load of the above-mentioned internal pressure can be implemented by, for example, sealing the pipe end with a sealing pad of rubber material and loading the pipe internal water pressure. In addition, in order to stabilize the shape, a mold of a desired diameter can also be used as an outer frame as required.

[0163] In addition, the plate thickness of the electric resistance welded steel pipe cited as an example of the steel pipe of the present invention is preferably 5 mm or more and preferably 30 mm or less.

[0164] In addition, in the sizing process, bending deformation occurs in the tube axis direction along the roller shape when the roller passes, and residual stress in the tube axis direction is generated. The greater the bending strain in the bending deformation, the greater the absolute value of the residual stress in the tube axis direction. The smaller the diameter of the sizing roller and the greater the thickness of the hot-rolled steel sheet, the greater the bending strain.

[0165] Therefore, in the present invention, from the viewpoint of reducing the shear residual stress, the diameter of the sizing roll is set to satisfy the above formula (1) in order to reduce the absolute value of the residual stress in the pipe axial direction.

[0166] When the diameter of the sizing roll is smaller than the right side of the above formula (1), the target shear residual stress in the present invention cannot be obtained. It should be noted that there is no particular upper limit on the diameter of the sizing roll, but if the sizing roll becomes larger, the load on the equipment increases, so the diameter of the sizing roll is preferably 2000 mm or less.

[0167] In the internal pressure loading process, the electric resistance welded steel tube material is expanded to generate tensile stress in the tube circumferential direction, thereby reducing the absolute value of the residual stress in the tube circumferential direction.

[0168] The greater the internal pressure p (MPa) in the internal pressure loading process, the smaller the absolute value of the residual stress in the pipe circumferential direction. The greater the radius of the steel pipe and the smaller the wall thickness of the steel pipe, the higher the tensile stress generated in the pipe circumferential direction.

[0169] The left side (X) of the above formula (2) corresponds to the internal pressure p when the tensile stress generated in the tube circumferential direction is equal to the yield stress of the electric resistance welded steel tube material.

[0170] In the present invention, from the viewpoint of reducing the shear residual stress, in order to reduce the absolute value of the residual stress in the tube axial direction, the internal pressure p is set to a value greater than the left side (X) of the formula (2), and the electric resistance welded steel tube billet is expanded to the plastic region. On the other hand, if the internal pressure p exceeds the right side (X×1.5) of the formula (2), the absolute value of the residual stress in the tube circumferential direction becomes smaller, but the amount of work hardening due to the expansion becomes too large, the dislocation density on the tube surface increases, and the fatigue properties in hydrogen decrease.

[0171] Under the above conditions, it is possible to obtain a steel material and a steel pipe having excellent fatigue properties in hydrogen that satisfy a predetermined crack growth rate in hydrogen.

[0172] Example 1

[0173] The following is a description of the embodiments that verify the effects of the present invention. It should be noted that in the following embodiments, steel materials and steel pipes were manufactured under the following manufacturing conditions, and the characteristics were evaluated. Steels with chemical compositions shown in Tables 1-1 and 1-2 were melted and cast into slabs. Steel materials No. 1 to 71 and 80 were all heated to 1200°C, and then hot rolled under the condition that the rolling end temperature was above 950°C. They were cooled by water cooling (accelerated cooling) to the cooling stop temperature shown in Tables 1-1 and 1-2 to manufacture steel materials. Steel pipes No. 2 to 71 and 80 were also formed into steel pipes using the obtained steel materials. No. 2 to 15, 33 to 44, 63 to 70, and 80 were subjected to the following pipe making process: the obtained steel (hot-rolled steel plate) was bent, the two ends were butted and welded; No. 16 to 32, 45 to 62, and 71 were subjected to the following pipe making process: the obtained steel (hot-rolled steel plate) was formed into a cylindrical shape by cold roll forming, the two circumferential ends of the cylindrical shape were butted and resistance welded; and the steel pipe was formed. The average cooling rate during cooling was within the range of 12°C / s±2°C / s. It should be noted that the temperature measurement of the steel was carried out using a thermocouple inserted into the center of the plate thickness. Furthermore, steel No. 1 to 31, 33 to 80 and steel pipe No. 2 to 31, 33 to 80 were placed at room temperature for more than 96 hours (h) for dehydrogenation treatment before evaluating the steel and steel pipe. Room temperature refers to 20±10°C. It should be noted that steel No. 32 was not subjected to dehydrogenation treatment. Steel materials No. 70, 71 and steel pipes No. 70, 71 cannot be welded and cannot be used as steel pipes, and therefore cannot be evaluated.

[0174] In addition, in steels No. 72 to 79, slabs having the same composition as that of steel No. 58 were used, and a study of manufacturing conditions was also conducted. In particular, for conditions not described, the steels No. 1 to 71 were manufactured under the same conditions as those described above. The heating temperature of steel No. 72 was 1300°C, so the heating temperature exceeded the upper limit of the invention. The rolling end temperature of steel No. 73 was 700°C, so the rolling end temperature was lower than the lower limit of the invention. Steel No. 74 was manufactured at a rolling end temperature of 900°C, which was within the scope of the invention. In addition, in steel No. 75, the average cooling rate after rolling was 9°C / s, which was outside the scope of the invention. In steel No. 76, the average cooling rate after rolling was 65°C / s, which was within the scope of the invention.

[0175] In Steel Material No. 77, when the heating temperature was 950° C., the slab became hard and could not be rolled, so the evaluation was not performed.

[0176] In addition, in steel material No. 78, the cooling stop temperature was 700° C., which is outside the scope of the present invention. In steel material No. 79, the cooling stop temperature was 240° C., which is outside the scope of the present invention.

[0177] The obtained steel materials were also used for Steel Materials No. 72 to 79, and were subjected to a pipe-making process of forming into a cylindrical shape by cold rolling, butting both circumferential ends of the cylindrical shape and performing electric resistance welding to obtain steel pipes.

[0178] The investigation of fatigue crack propagation characteristics was evaluated by fatigue crack growth test. From each steel material, a CT (compact tension) test piece (a nearly square test piece with a notch on one end) according to ASTM E 647 was taken in a manner that the load direction was parallel to the rolling direction. The length of the fatigue crack was measured by the compliance method using a clamp gauge, and the fatigue crack propagation rate in 5MPa high-pressure hydrogen was calculated. It should be noted that for the test piece, when the plate thickness is less than 10mm, it is ground from the surface every 0.5mm to make 2mm, 5mm, 8mm, and 9mm respectively. In the case of plate thicknesses other than these, a 10mm thick test piece is taken from the position of t / 2 (t: plate thickness), and mirror polishing is performed on the front and back of the crack extension part. At this time, as the stable growth area where the Paris law is established, the stress expansion coefficient range ΔK=20 (MPa·m 1 / 2 The fatigue crack growth rate (m / cycle) in the test specimen was evaluated as a representative value. The results are shown in Table 1. The crack growth rate da / dN was 1.0×10 － 6 m·cycle -1 The following levels are considered acceptable.

[0179] In addition, Nb precipitates were evaluated by the following method. Square steels randomly taken from three locations of each sample were machined to a mirror surface and then etched. The size was not specifically specified, but the surface area was uniformly 10 cm 2 Next, the Nb precipitates were observed and analyzed by EPMA. 10 / μm 2 The case of Nb precipitates of 2 nm to 100 nm or more is designated as Y, and the case of Nb precipitates smaller than the above is designated as N, and are described in Tables 1-1 and 1-2.

[0180] It should be noted that the same results as those of the steel materials were obtained for the steel pipes in Tables 1-1, 1-2, and 2.

[0181] All the examples of the present invention satisfy the crack growth rate da / dN in hydrogen of 1.0×10 －6 m / cycle or less.

[0182] [Table 1-1]

[0183]

[0184] [Table 1-2]

[0185]

[0186] Example 2

[0187] The following is a description of the examples that verified the effects of the present invention. It should be noted that in the following examples, steels were manufactured under the following manufacturing conditions and their characteristics were evaluated. Using slabs having the same composition as steel No. 17, 26, 55, and 67 shown in Tables 1-1 and 1-2, the steels were manufactured under the same conditions as steel No. 17, 26, 55, and 67 shown in Example 1 until the controlled cooling process, to obtain steel No. 17A, 26A, 55A, and 67A. Steel No. 17A, 26A, and 55A were subjected to the following pipe making process: the obtained steel (hot-rolled steel plate) was formed into a cylindrical shape by cold rolling, and the circumferential ends of the cylindrical shape were butted and resistance welded; steel 67A was subjected to the following pipe making process: the obtained steel (hot-rolled steel plate) was bent, the two ends were butted and welded, and steel pipes were formed. As shown in Table 2, the characteristics were evaluated when the dehydrogenation treatment conditions were changed. The results are shown in Table 2.

[0188] The dehydrogenation treatment of steel materials No. 17, 26, 55, and 67 in Example 1 was carried out with the dehydrogenation treatment temperature T (atmosphere temperature) at room temperature and the holding time being 96 hours, but in this example, the dehydrogenation treatment temperature T (atmosphere temperature) was all carried out at 50°C. Steel pipes No. 81, 84, 86, and 88 were carried out in such a manner that the dehydrogenation treatment temperature T (atmosphere temperature) was 50°C and the holding time tc after the plate thickness center temperature Tc reached 50°C satisfied the formula (A). Steel pipes No. 82, 85, 87, and 89 were carried out in such a manner that the dehydrogenation treatment temperature T (atmosphere temperature) was 50°C and the holding time t at the dehydrogenation treatment temperature T was 50°C satisfied the above formula (A), but the holding time tc after the plate thickness center temperature Tc reached 50°C did not satisfy the above formula (A).

[0189] In steel pipe No. 83, the dehydrogenation treatment temperature T (atmosphere temperature) is 50°C, but the holding time t of the atmosphere temperature and the holding time tc after the plate center temperature Tc reaches 50°C do not satisfy the above formula (A).

[0190] In Table 2, "dehydrogenation holding time t is Y" means that the dehydrogenation treatment temperature T (atmosphere temperature) is 50°C and the holding time t satisfies the formula (A), and "dehydrogenation holding time t is N" means that the dehydrogenation treatment temperature T (atmosphere temperature) is 50°C, but the holding time t does not satisfy the formula (A). In addition, "holding time tc of the steel material center temperature Tc is Y" means that the holding time tc after the plate thickness center temperature Tc reaches 50°C satisfies the formula (A), and "holding time tc of the steel material center temperature Tc is N" means that the plate thickness center temperature Tc reaches 50°C, but the holding time tc after Tc reaches 50°C does not satisfy the formula (A).

[0191] All the examples of the present invention satisfy the crack growth rate da / dN in hydrogen of 1.0×10 －6 Among them, when the dehydrogenation treatment is carried out under more preferred conditions, the crack propagation characteristics are more excellent.

[0192] [Table 2]

[0193]

Claims

1. A steel material having excellent fatigue properties in hydrogen, comprising, by mass%, C: 0.02-0.15%, Si: 0.05-0.5%, Mn: 0.3-2.0%, Al: 0.01-0.15%, N: 0.0005-0.008%, P: 0.03% or less, S: 0.01% or less, O: 0.01% or less, Nb: more than 0% and 2.5% or less, H: 0.0010% The following, Cu: 0-2.5%, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.03%, Sb: 0-0.3%, Ca: 0-0.01%, Mg: 0-0.01%, REM: 0-0.005%, the remainder is composed of Fe and inevitable impurities; Nb precipitates with equivalent circle diameter of 2nm to 100nm are 10 / μm 2 above, The crack growth rate da / dN when the stress expansion coefficient in hydrogen above 1MPa = 20MPa√m is 1.0×10 －6 m·cycle -1 the following.

2. A method for manufacturing a steel material, comprising: A heating step of heating the slab having the component composition of claim 1 at 1000 to 1250° C. In the hot rolling step, the slab heated in the heating step is rolled at a finishing temperature of Ar3 point or higher. a controlled cooling step of cooling the steel plate obtained in the hot rolling step at an average cooling rate of 10°C / s or more at a temperature at the center of the plate thickness of 1000°C to 400°C and a cooling stop temperature of 250°C to 650°C; and The dehydrogenation treatment step is to perform a dehydrogenation treatment on the steel sheet obtained in the controlled cooling step.

3. A steel pipe having excellent fatigue properties in hydrogen, comprising, by mass%, C: 0.02-0.15%, Si: 0.05-0.5%, Mn: 0.3-2.0%, Al: 0.01-0.15%, N: 0.0005-0.008%, P: 0.03% or less, S: 0.01% or less, O: 0.01% or less, Nb: more than 0% and 2.5% or less, H: 0.0010% The following, Cu: 0-2.5%, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.03%, Sb: 0-0.3%, Ca: 0-0.01%, Mg: 0-0.01%, REM: 0-0.005%, the remainder is composed of Fe and inevitable impurities; Nb precipitates with equivalent circle diameter of 2nm to 100nm are 10 / μm 2 above, The crack growth rate da / dN when the stress expansion coefficient in hydrogen above 1MPa = 20MPa√m is 1.0×10 －6 m·cycle -1 the following.

4. A method for manufacturing a steel pipe, comprising: A heating step of heating the slab having the component composition of claim 3 at 1000 to 1250° C. In the hot rolling step, the slab heated in the heating step is rolled at a finishing temperature of Ar3 point or higher. A controlled cooling step is performed to cool the hot rolled steel sheet obtained in the hot rolling step at an average cooling rate of 10°C / s or more at a temperature at the center of the sheet thickness at 1000°C to 400°C and a cooling stop temperature of 250°C to 650°C. Any of the following pipe making processes: a pipe making process in which, after the controlled cooling process, the hot rolled steel sheet is bent, both ends are butted and welded; and a pipe making process in which, after the controlled cooling process, the hot rolled steel sheet is formed into a cylindrical shape by cold roll forming, both circumferential ends of the cylindrical shape are butted and resistance welded, and The dehydrogenation treatment step is to perform a dehydrogenation treatment on the steel pipe obtained in the pipe making step.

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