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

Abstract

The purpose of the present invention is to provide natural gas suitable as 100% hydrogen or hydrogen having a hydrogen partial pressure of 1 MPa or more (natural gas being gas mainly composed of hydrocarbons such as methane or ethane) A steel pipe having excellent fatigue characteristics in hydrogen in a high-pressure hydrogen environment, which is used in a steel structure used in a high-pressure hydrogen environment, such as a line pipe, a method for manufacturing the steel pipe, a steel material, and a method for manufacturing the steel material. A steel pipe having excellent fatigue characteristics in hydrogen, the steel pipe 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 at a stress expansion coefficient range of 20 MPa square root of m in hydrogen at 1 MPa or more.

Description

Technical Field

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

[0002] As an existing energy infrastructure, there are pipelines for natural gas transmission. These steel materials are required to suppress the occurrence of hydrogen-induced cracking in sour environments. On the other hand, in recent years, hydrogen has attracted great attention worldwide as a clean energy source 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 is pressurized and sent through natural gas pipelines, and hydrogen is used as an alternative. 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 materials used in such an environment will suffer from "hydrogen embrittlement" in which hydrogen intrudes into the steel and deteriorates its properties. Therefore, it is necessary to have not only the high toughness and sour resistance required of conventional pipelines, but also the resistance to hydrogen embrittlement required in a hydrogen environment.

[0003] Steel structures used in high-pressure hydrogen environments have traditionally utilized austenitic stainless steels such as SUS316L, which are less susceptible to hydrogen embrittlement than low-alloy steels. However, austenitic stainless steels such as SUS316L are not only expensive but also have low strength. Therefore, designing them to withstand high hydrogen pressures requires thicker walls, which increases the cost of the hydrogen-resistant structures themselves. Therefore, there is a strong demand for low-alloy steels that are more cost-effective and can withstand high-pressure hydrogen environments as steel materials for hydrogen-resistant steel structures.

[0004] To meet such demands, for example, the steel for high-pressure hydrogen environment described in Patent Document 1 is steel used in high-pressure hydrogen environment. By making Ca / S less than 1.5 or greater than 11, the diffusible hydrogen concentration ratio is reduced, thereby suppressing embrittlement caused by diffusible hydrogen.

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

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

[0007] Furthermore, Patent Document 4 proposes a low-alloy steel for use in high-pressure hydrogen environments. The low-alloy steel described in Patent Document 4 further increases the Mo content compared to existing steels by adding V, raising the tempering temperature, and utilizing V-Mo carbides. This improves the carbide morphology at the grain boundaries and significantly enhances resistance to hydrogen environment embrittlement.

[0008] Furthermore, Patent Document 5 proposes steel for high-pressure hydrogen storage containers with excellent hydrogen resistance. According to the technology described in Patent Document 5, by performing a long stress relief annealing after normalizing during steel plate production, MC carbides (Mo, V)C are dispersed and precipitated finely and densely, thereby improving the steel's hydrogen resistance, including resistance to hydrogen embrittlement.

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

[0010] Prior art literature

[0011] Patent Literature

[0012] Patent Document 1: Japanese Patent Application Laid-Open No. 2005-2386

[0013] Patent Document 2: Japanese Patent Application Laid-Open No. 2009-46737

[0014] Patent Document 3: Japanese Patent Application Laid-Open No. 2009-275249

[0015] Patent Document 4: Japanese Patent Application Laid-Open No. 2009-74122

[0016] Patent Document 5: Japanese Patent Application Laid-Open No. 2010-37655

[0017] Patent Document 6: Japanese Patent Application Laid-Open No. 2012-107332

[0018] Non-patent literature

[0019] Non-Patent Document 1: Matsunaga et al., Int J Hydrogen Energy, Vol. 40 (2015), pp. 5739-5748 Summary of the Invention

[0020] Fluctuations in pressure within line pipes during operation and periodic closures repeatedly stress structures like line pipes. Therefore, fatigue fracture must be considered when designing steel structures like line pipes. However, as shown in Non-Patent Document 1, it is known that the fatigue life of materials decreases in high-pressure hydrogen environments. This means that when designing line pipes based on conventional natural gas line pipes, the service life of the pipes is reduced. However, while these prior arts can suppress hydrogen-induced cracking in acidic environments, they cannot sufficiently improve fatigue strength in hydrogen. Specifically, they face the problem of difficulty in achieving high fatigue strength in hydrogen, which is more likely to affect service life, in addition to suppressing hydrogen-induced cracking in acidic environments.

[0021] In view of the above-mentioned problems of the prior art, the present invention aims to provide a steel pipe having excellent fatigue properties in hydrogen in a high-pressure hydrogen environment, a method for manufacturing the same, a steel material, and a method for manufacturing the same, which are suitable for use in steel structures used in high-pressure hydrogen environments, such as line pipes for natural gas (natural gas is a gas mainly composed of hydrocarbons such as methane and ethane) that is 100% hydrogen or contains hydrogen at a hydrogen partial pressure of 1 MPa or more.

[0022] It should be noted that the term "excellent fatigue properties in hydrogen under high-pressure hydrogen" herein refers to a crack growth rate da / dN of 1.0×10-1 at a stress enlargement coefficient range = 20 MPa√m obtained by fatigue testing in accordance with ASTM E647 at room temperature (20±10°C) in a hydrogen atmosphere at a pressure of 1 MPa or higher or a natural gas (mainly composed of hydrocarbons such as methane and ethane) containing hydrogen at a hydrogen partial pressure of 1 MPa or higher, with a stress ratio of R = 0.1 and a stress enlargement coefficient range = 20 MPa√m = 1.0×10-1 －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 by volume fraction and a total gas pressure of 30 MPa or less.

[0023] It should be noted that if the crack growth rate da / dN in 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.

[0024] From the above viewpoints, the present inventors have conducted intensive research on the conditions that various steel materials should satisfy in hydrogen, and have discovered a new steel pipe and steel material having excellent fatigue properties in hydrogen.

[0025] The present invention has been made based on the above-mentioned new findings and further studies, and its gist is as follows.

[0026] [1] A steel pipe having excellent fatigue properties in hydrogen, comprising, in mass%, C: 0.10-0.45%, Si: 0.01-2.0%, Mn: 0.3-2.0%, Al: 0.01-0.15%, N: 0.0005-0.008%, P: 0.015% or less, S: 0.0015% or less, O: 0.01% or less, H: 0.0010% or less, and Cu: 0-2. 5%, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, Nb: 0-0.5%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.3%, 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;

[0027] Retained austenite is less than 3%,

[0028] The number of carbides with a diameter of 200 nm or more is 20 per 10 μm. 2 the following,

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

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

[0031] The casting step is to cast the steel billet having the composition described in [1] at a casting speed of 1.8 m / min or less.

[0032] The heating process is carried out at a temperature below 1350°C.

[0033] In the hot rolling process, the steel billet heated in the above heating process is rolled under the condition of a rolling end temperature of 820°C or above to form a steel pipe shape.

[0034] In the cooling step, the steel pipe obtained in the hot rolling step is kept at a temperature of Ac3 point to 1000°C, and then cooled under the following conditions: Group A or Group B, and

[0035] In the tempering step, the steel pipe obtained in the cooling step is tempered at 400° C. to Ac1 point for less than 60 minutes.

[0036] Group A:

[0037] The steel pipe is cooled to 50°C or lower at an average cooling rate of 15°C / s or higher at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe from 800°C to 550°C, and at an average cooling rate of 15°C / s or lower at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe from 550°C to 50°C.

[0038] Group B:

[0039] The steel pipe is cooled to 50°C or lower at an average cooling rate from 800°C to 300°C of 10°C / s or higher at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe, and at an average cooling rate from 300°C to 50°C of 5°C / s or lower at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe.

[0040] [3] The method for manufacturing a steel pipe according to [2], wherein a quenching step is provided before the tempering step, wherein the steel pipe is reheated to a temperature of Ac3 point to 1000°C, and the cooling conditions are the following Group A or Group B.

[0041] Group A:

[0042] The steel pipe is cooled to 50°C or lower at an average cooling rate of 15°C / s or higher at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe from 800°C to 550°C, and at an average cooling rate of 15°C / s or lower at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe from 550°C to 50°C.

[0043] Group B:

[0044] The steel pipe is cooled to 50°C or lower at an average cooling rate from 800°C to 300°C of 10°C / s or higher at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe, and at an average cooling rate from 300°C to 50°C of 5°C / s or lower at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe.

[0045] [4] The method for manufacturing a steel pipe according to [2] or [3], wherein the casting speed is 1.0 m / min or less.

[0046] [5] A steel material having excellent fatigue properties in hydrogen, comprising, in mass%, C: 0.10-0.45%, Si: 0.01-2.0%, Mn: 0.3-2.0%, Al: 0.01-0.15%, N: 0.0005-0.008%, P: 0.015% or less, S: 0.0015% or less, O: 0.01% or less, H: 0.0010% or less, and Cu: 0-2. 5%, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, Nb: 0-0.5%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.3%, 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;

[0047] Retained austenite is less than 3%,

[0048] The number of carbides with a diameter of 200 nm or more is 20 per 10 μm. 2 the following,

[0049] The stress expansion coefficient range in hydrogen above 1 MPa = crack growth rate da / dN when 20 MPa√m is 1.0×10 －6 m·cycle -1 the following.

[0050] [6] A method for manufacturing a steel material, comprising:

[0051] The casting step is to cast the steel billet having the composition described in [5] at a casting speed of 1.8 m / min or less.

[0052] The heating process is carried out at a temperature below 1350°C.

[0053] In the hot rolling process, the steel billet heated in the above heating process is rolled at a rolling end temperature of 820°C or above.

[0054] In the cooling step, the steel obtained in the hot rolling step is kept at a temperature of Ac3 point to 1000°C, and then cooled under the following conditions: Group A or Group B, and

[0055] In the tempering step, the steel material obtained in the cooling step is tempered at 400° C. to Ac1 point for less than 60 minutes.

[0056] Group A:

[0057] The steel is cooled to 50°C or below at an average cooling rate of 15°C / s or above at a position 1 / 4 of the wall thickness from the steel surface from 800°C to 550°C, and at an average cooling rate of 15°C / s or below at a position 1 / 4 of the wall thickness from the steel surface from 550°C to 50°C.

[0058] Group B:

[0059] The steel is cooled to 50°C or below at an average cooling rate of 10°C / s or more from 800°C to 300°C at a position 1 / 4 of the wall thickness from the steel surface, and at an average cooling rate of 5°C / s or less from 300°C to 50°C at a position 1 / 4 of the wall thickness from the steel surface.

[0060] [7] The method for manufacturing a steel material according to [6], wherein a quenching step is provided before the tempering step, wherein the steel material is reheated to a temperature of Ac3 point to 1000°C, and the cooling conditions are the following group A or group B.

[0061] Group A:

[0062] The steel is cooled to 50°C or below at an average cooling rate of 15°C / s or above at a position 1 / 4 of the wall thickness from the steel surface from 800°C to 550°C, and at an average cooling rate of 15°C / s or below at a position 1 / 4 of the wall thickness from the steel surface from 550°C to 50°C.

[0063] Group B:

[0064] The steel is cooled to 50°C or below at an average cooling rate of 10°C / s or more from 800°C to 300°C at a position 1 / 4 of the wall thickness from the steel surface, and at an average cooling rate of 5°C / s or less from 300°C to 50°C at a position 1 / 4 of the wall thickness from the steel surface.

[0065] [8] The method for manufacturing a steel material according to [6] or [7], wherein the casting speed is 1.0 m / min or less.

[0066] According to the present invention, a steel pipe and a steel material having extremely excellent fatigue properties in a high-pressure hydrogen environment can be obtained, and the present invention is extremely useful industrially. DETAILED DESCRIPTION

[0067] Next, the method for implementing the present invention will be described in detail.

[0068] As a first embodiment, a method for implementing a steel pipe will be specifically described. Next, as a second embodiment, a method for implementing a steel material will be specifically described.

[0069] First embodiment

[0070] [Ingredients]

[0071] The reasons for limiting the chemical composition of the steel pipe (including steel materials) of the present invention are described below. It should be noted that "%" in the following description means "mass %" unless otherwise specified.

[0072] C: 0.10~0.45%

[0073] C is an element necessary for increasing strength. Its effect is insufficient when the content is less than 0.10%. Therefore, the C content is 0.10% or more. The C content is preferably 0.13% or more. The C content is more preferably 0.15% or more, and even more preferably 0.18% or more. On the other hand, if the C content exceeds 0.45%, quenching cracks may sometimes occur during quenching, and it may also cause the formation of coarse carbides, deteriorating fatigue properties in hydrogen. Therefore, the C content is 0.45% or less. The C content is preferably 0.43% or less. The C content is more preferably 0.40% or less, and even more preferably 0.38% or less.

[0074] Si: 0.01-2.0%

[0075] Si is contained as a deoxidizing agent during the steelmaking stage and as an element that ensures hardenability. If the content is less than 0.01%, its effect is insufficient, so the Si content is set to 0.01% 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 the content exceeds 2.0%, grain boundaries become embrittled, deteriorating low-temperature toughness and fatigue properties in hydrogen. Therefore, the Si content is set to 2.0% or less. The Si content is preferably 1.5% or less. The Si content is preferably 1.0% or less, and more preferably 0.8% or less.

[0076] Mn: 0.3-2.0%

[0077] Mn is contained as an element to ensure hardenability. When it is less than 0.3%, its effect is not sufficient, so the Mn content 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. Furthermore, if there is too much Mn, the austenite stability increases, so it is possible to exceed the specified amount of retained austenite and increase the amount of hydrogen in the steel. In addition, the hardness of the surface part and the central segregation part increases during cooling (accelerated cooling or quenching) after hot rolling, so the fatigue properties in hydrogen deteriorate. Therefore, the Mn content is 2.0% or less. The Mn content is more preferably 1.5% or less, and the Mn content is further preferably 1.3% or less. The Mn content is most preferably 1.0% or less.

[0078] Al: 0.01~0.15%

[0079] Al is contained as a deoxidizing material and also has the effect of pinning austenite grains during heating as fine precipitates of Al-based nitrides, thereby suppressing grain coarsening. However, this effect is insufficient when the content is less than 0.01%. 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 exceeds 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 0.13% or less. The Al content is more preferably 0.10% or less, and even more preferably 0.08% or less.

[0080] N: 0.0005~0.008%

[0081] N is contained in order to form fine precipitates by forming nitrides with Nb, Ti, Al, etc., and to suppress grain coarsening and improve low-temperature toughness by pinning austenite grains during heating. When the content is less than 0.0005%, the microstructure cannot be fully refined, so the N content is set to 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 the content exceeds 0.008%, the amount of dissolved N increases, which impairs the toughness of the base material and the weld heat-affected zone, and the fatigue properties in hydrogen deteriorate. Therefore, the N content is set to 0.008% or less. The N content is preferably 0.007% or less. The N content is more preferably 0.006% or less, and even more preferably 0.005% or less.

[0082] P: 0.015% or less

[0083] P, an impurity element, tends to segregate at grain boundaries. If it exceeds 0.015%, it reduces the bonding strength between adjacent grains and degrades low-temperature toughness and fatigue properties in hydrogen. Therefore, the P content is 0.015% or less. The P content is preferably 0.013% or less, and more preferably 0.010% or less. While the lower limit is not particularly limited, it is preferably 0.001% or more to avoid increased costs.

[0084] S: 0.0015% or less

[0085] As an impurity element, S tends to segregate at grain boundaries and easily form MnS as a non-metallic inclusion. If the S content exceeds 0.0015%, the bonding strength between adjacent grains decreases, the amount of inclusions increases, and low-temperature toughness and fatigue properties in hydrogen deteriorate. Therefore, the S content is 0.0015% or less. The S content is preferably 0.0013% or less. The S content is more preferably 0.0010% or less, and even more preferably 0.0008% or less. There is no particular lower limit, but from the perspective of increased cost, it is preferably 0.0001% or more.

[0086] O: 0.01% or less

[0087] O affects the workability of the material by forming oxides with Al and other materials, so the lower the content, the better. When the content exceeds 0.01%, inclusions increase, impairing workability. Furthermore, as inclusions increase, fatigue properties in hydrogen deteriorate. Therefore, the O content is 0.01% or less. The O content is preferably 0.009% or less. The O content is more preferably 0.008% or less. There is no particular lower limit, but from the perspective of increasing costs, it is preferably 0.0001% or more. The O content is more preferably 0.002% or more.

[0088] H: 0.0010% or less

[0089] H is sometimes introduced into steel during various manufacturing processes. 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, the H content is less than 0.0002%. On the other hand, if it is less than 0.00001%, it becomes the main reason for the increase in cost, so it is preferably more than 0.00001%. The H content is 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.

[0090] In the present invention, the remainder of the above-mentioned composition is preferably a steel composition consisting of Fe and inevitable impurities, but depending on the desired characteristics, it is preferred to further appropriately contain one or more of Cu: 0-2.5%, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, Nb: 0-0.5%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.3%, Sb: 0-0.3%, Ca: 0-0.01%, Mg: 0-0.01%, and REM: 0-0.005%.

[0091] Cu: 0-2.5%

[0092] Cu has the effect of improving hardenability. Therefore, when Cu is included, the Cu content can be 0% or more. However, if it is less than 0.05%, the above effect is difficult to achieve, so the Cu content is preferably 0.05% or more. On the other hand, if it exceeds 2.5%, hot cracks are likely to occur when the steel sheet is heated. Therefore, when Cu is included, 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 even more preferably 1.8% or less.

[0093] Ni: 0-2.5%

[0094] Ni, like Cu, has the effect of improving hardenability and further improves toughness. Therefore, when Ni is included, the Ni content can be 0% or more. However, if it is less than 0.05%, it is difficult to achieve the above-mentioned effect, so the Ni content is preferably 0.05% or more. On the other hand, if it exceeds 2.5%, the economic efficiency is poor. Therefore, when Ni is included, it is 2.5% or less. The Ni content is preferably 2.3% or less. More preferably, it is 2.0% or less, and preferably 1.8% or less.

[0095] Cr: 0~2.5%

[0096] Cr is contained as an element to ensure hardenability. When Cr is contained, the Cr content can be 0% or more. However, if it is less than 0.1%, the above-mentioned effect is difficult to achieve. Therefore, the Cr content is preferably 0.1% or more. On the other hand, if it exceeds 2.5%, toughness deteriorates and economic efficiency is poor. Therefore, when Cr is contained, it 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 1.8% or less, and most preferably 1.5% or less.

[0097] Mo: 0-2.0%

[0098] Mo has the effect of improving hardenability. Therefore, when Mo is included, the Mo content can be 0% or more. However, if it is less than 0.05%, the above effect is difficult to achieve. Therefore, the Mo content is preferably 0.05% or more. On the other hand, if it exceeds 2.0%, the economic efficiency is poor. Therefore, when Mo is included, it 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 even more preferably 1.2% or less.

[0099] Nb: 0-0.5%

[0100] Nb improves hardenability and, as fine precipitates of Nb-based carbonitrides, pins austenite grains during heating, suppressing grain coarsening. Therefore, when Nb is included, the Nb content can be 0% or more, but if it is less than 0.005%, the above-mentioned effect is difficult to achieve, so the Nb content is preferably 0.005% or more. The Nb content is more preferably 0.01% or more. On the other hand, if the content exceeds 0.5%, there is a concern that coarse Nb carbonitrides will precipitate, leading to deterioration in toughness. Therefore, when Nb is included, the Nb content is 0.5% or less. The Nb content is preferably 0.4% or less. The Nb content is preferably 0.3% or less, and the Nb content is preferably 0.2% or less.

[0101] V: 0~0.5%

[0102] V has the effect of improving hardenability, and as fine precipitates of V-based carbides, it pins austenite grains during heating, inhibiting grain coarsening. 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 achieve the above-mentioned effect, so the V content is preferably 0.005% or more. On the other hand, when the content exceeds 0.5%, there is a concern that coarse V carbonitrides will precipitate, leading to deterioration of toughness. 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 even more preferably 0.2% or less.

[0103] Ti: 0~0.5%

[0104] Ti has the effect of improving hardenability and has the effect of pinning austenite grains during heating as fine precipitates of Ti-based carbonitrides, thereby inhibiting grain growth. Therefore, when Ti is contained, the Ti content can be 0% or more, but when it is less than 0.005%, it is difficult to achieve the above-mentioned effect, so the Ti content is preferably 0.005% or more. The Ti content is preferably 0.01% or more. On the other hand, when the content exceeds 0.5%, coarse angular nitrides are easily formed, and toughness 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 even more preferably 0.2% or less.

[0105] W: 0~2.5%

[0106] W has the effect of improving hardenability. Therefore, when W is included, the W content can be 0% or more. However, if it is less than 0.05%, this effect is difficult to achieve, so the W content is preferably 0.05% or more. On the other hand, if it exceeds 2.5%, economic efficiency is poor. Therefore, when W is included, 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 even more preferably 1.8% or less.

[0107] B: 0~0.005%

[0108] B is an element that ensures hardenability. Therefore, when B is included, the B content can be 0% or more. However, if it is less than 0.0005%, the above effect is difficult to achieve. Therefore, the B content is preferably 0.0005% or more. On the other hand, if it exceeds 0.005%, the toughness deteriorates. Therefore, when B is included, 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 even more preferably 0.002% or less.

[0109] Sn: 0~0.3%

[0110] Sn improves the corrosion resistance of steel pipes. Therefore, when Sn is included, the Sn content can be 0% or higher. However, if it is less than 0.005%, the aforementioned effect is difficult to achieve. Therefore, the Sn content is preferably 0.005% or higher. The Sn content is more preferably 0.01% or higher. On the other hand, if the Sn content exceeds 0.3%, high-temperature ductility decreases, increasing the likelihood of cracking during casting. Therefore, when Sn is included, the Sn content is 0.3% or lower. The Sn content is preferably 0.25% or lower. The Sn content is more preferably 0.2% or lower, and even more preferably 0.15% or lower.

[0111] Sb: 0~0.3%

[0112] Sb improves the corrosion resistance of steel pipes. Therefore, when Sb is included, the Sb content can be 0% or more. However, if it is less than 0.005%, the aforementioned effect is difficult to achieve. Therefore, the Sb content is preferably 0.005% or more. The Sb content is more preferably 0.01% or more. On the other hand, if the Sb content exceeds 0.3%, high-temperature ductility decreases, and hot rolling properties are also impaired. Therefore, when Sb is included, 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 even more preferably 0.15% or less.

[0113] Ca: 0~0.01%

[0114] Ca forms CaS, which has the effect of controlling the morphology of sulfide inclusions, replacing MnS, which is an inclusion that easily elongates due to rolling, with spherical inclusions, i.e., CaS, that are less susceptible to elongation due to rolling. Therefore, when Ca is included, the Ca content can be 0% or more, but if it is less than 0.0005%, the above effect is difficult to achieve, so the Ca content is preferably 0.0005% or more. The Ca content is more preferably 0.001% or more. On the other hand, if the content exceeds 0.01%, cleanliness is reduced, and thus the material properties such as toughness deteriorate. Therefore, when Ca is included, 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 even more preferably 0.002% or less.

[0115] Mg: 0~0.01%

[0116] Mg is sometimes used as a desulfurizing agent for molten iron. Therefore, when Mg is included, the Mg content can be 0% or higher. However, if it is less than 0.0005%, the aforementioned effects are difficult to achieve. Therefore, the Mg content is preferably 0.0005% or higher. The Mg content is more preferably 0.001% or higher. On the other hand, a content exceeding 0.01% results in a decrease in cleanliness. Therefore, when Mg is included, the Mg content is 0.01% or lower. The Mg content is preferably 0.005% or lower. The Mg content is more preferably 0.004% or lower, and even more preferably 0.003% or lower.

[0117] REM: 0~0.005%

[0118] REM reduces the amount of dissolved S at grain boundaries and improves SR crack resistance by generating sulfides as REM (O, S) in steel. Therefore, when REM is contained, the REM content can be 0% or more, but when it is less than 0.0005%, the above effect is difficult to obtain, so the REM content is preferably 0.0005% or more. On the other hand, when the content exceeds 0.005%, REM sulfides accumulate significantly in the precipitation bands, leading to material degradation. 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.

[0119] In the chemical compositions of the steel plates and steel pipes, the remainder other than the above-mentioned components (elements) is composed of Fe and inevitable impurity elements.

[0120] The preferred metal structure of the steel pipe of the present invention will be described in detail.

[0121] Retained austenite is less than 3%

[0122] Residual austenite in steel pipes can increase the amount of hydrogen in the steel, increasing susceptibility to hydrogen embrittlement. Furthermore, when austenite transforms to martensite due to stress loads during use, hydrogen-induced cracking can easily occur due to the extremely hard nature of martensite, sometimes initiating cracks in the martensite. In the present invention, the fatigue crack growth rate is reduced by limiting the retained austenite content to 3% or less. It is preferably 2% or less, and more preferably 1% or less. The retained austenite content can be 0%.

[0123] The number of carbides with a diameter of 200 nm or more is 20 per 10 μm. 2 the following

[0124] In the steel pipe of the present invention, if coarsened carbides are present that are larger than a specified value, fatigue properties in hydrogen are adversely affected. Therefore, the number of carbides with a diameter of 200 nm or more is 20 per 10 μm. 2 Less than 15 pieces / 10μm. 2 Less than 10 pieces / 10 μm. More preferably, 10 pieces / 10 μm 2 Below, more preferably 5 / 10 μm 2 For the lower limit, the less the better, and it can be 0 / 10μm. 2 Here, the diameter is the value calculated from 2√(A / 2×B / 2) using the long side A and the short side B passing through the center. Furthermore, carbide refers to intermetallic compounds such as cementite, ε-carbide, χ-carbide, and Fe7C3.

[0125] Furthermore, when carbides with a diameter of less than 200 nm are uniformly dispersed in the grains at intervals of 100 nm or more, this will not adversely affect fatigue properties in hydrogen and will contribute to strength improvement. Therefore, the number of carbides with a diameter of less than 200 nm dispersed in the grains is preferably 10 / 10 μm. 2 As for the upper limit, if the precipitates are too large, the coarsening will proceed. For this reason, the number of carbides with a diameter of less than 200 nm dispersed in the grains is preferably 100 / 10 μm. 2 the following.

[0126] Note that carbide precipitation is affected by the conditions of the cooling step and tempering step after hot rolling, which will be described later. Therefore, it is important to control these conditions.

[0127] The crack growth rate da / dN when the stress expansion coefficient in hydrogen above 1 MPa is 20 MPa√m is 1.0×10 －6 m·cycle -1 the following

[0128] The fatigue crack growth rate is an important parameter in the design of steel pipes used in pipelines and gas containers. It is necessary to obtain a service life that ensures the safety of fractured structural components. In fractured structural components, it is difficult to make the cracking and crack generation position zero. When subjected to repeated stress, cracks will inevitably generate and grow. The crack growth rate is small when the stress state applied to the crack front is small, and it becomes larger when the stress state at the crack front increases. In a hydrogen environment, hydrogen penetrates into the steel pipe, making the crack easy to grow. 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 frequency: 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 range = 20MPa√m is 1.0×10 －6 m·cycle -1 Below, the service life of the steel structure in a high-pressure hydrogen environment can be fully ensured. Therefore, in the crack growth test in hydrogen above 1 MPa, the crack growth rate da / dN when the stress expansion coefficient is 20 MPa√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

[0129] －6 m·cycle -1 As for the lower limit, it is believed that the closer to the atmosphere, the better the result. The crack growth rate da / dN when the stress expansion coefficient is 20 MPa√m is more preferably 0.05×10 －6 m·cycle -1 above.

[0130] The thickness of the steel pipe is not particularly limited, but is preferably 5 mm or more and 30 mm or less.

[0131] Incidentally, examples of the steel pipe of the present invention include seamless steel pipes, electric resistance welded pipes, and UOE steel pipes. The method for producing a seamless steel pipe will be specifically described below as an example.

[0132] The steel pipe having excellent fatigue properties in hydrogen of the present invention only needs to have the above-mentioned component composition and satisfy the crack growth rate in hydrogen. The production method thereof will be described in detail below.

[0133] Of course, electric resistance welded pipes and UOE steel pipes can be manufactured by performing treatments so as to have the same thermal history.

[0134] The steel pipe of the present invention can be produced by sequentially performing the following steps (1) to (3).

[0135] (1) The process of casting after adjusting the composition of the steel billet

[0136] (2) The process of heating and rolling the cast material to obtain the shape of a steel pipe, hot rolling, and cooling (accelerated cooling) (including the case of reheating and quenching before the tempering process),

[0137] (3) A step of tempering the steel pipe obtained in the above steps.

[0138] The following describes each process. It should be noted that, unless otherwise specified, the temperatures described below refer to the temperature at the center of the steel billet or steel pipe's thickness. The average cooling rate refers to the temperature at a point 1 / 4 of the wall thickness from the inner surface of the steel pipe. The temperatures at the center of the wall thickness and at a point 1 / 4 of the wall thickness from the inner surface of the steel pipe are estimated from the steel pipe surface temperature measured with a radiation thermometer using heat transfer calculations that take into account the heat transfer coefficient of the steel material.

[0139] [Casting process]

[0140] Casting speed: 1.8m / min. or less

[0141] The slower the casting speed, the more it reduces the hydrogen concentration and inclusions in the steel. This effect becomes significant at speeds of 1.8 m / min or less. Therefore, the casting speed is set to 1.8 m / min or less. It is preferably 1.5 m / min or less. More preferably, it is 1.0 m / min or less. Even more preferably, it is 0.5 m / min or less. Most preferably, it is 0.1 m / min or less. There is no particular lower limit; the casting speed can be any speed greater than 0 m / min.

[0142] Heating process

[0143] For hot rolling, the steel slab having the above-mentioned composition is heated. The steel slab is not particularly limited, and for example, a slab or billet obtained by a conventional continuous casting method can be used.

[0144] Heating temperature: below 1350℃

[0145] If the heating temperature in the heating process exceeds 1350°C, the average grain size of the original austenite grains becomes too large and the various properties deteriorate, so the heating temperature is 1350°C or less. The heating temperature is more preferably 1300°C or less, further preferably 1250°C or less, and most preferably 1200°C or less. On the other hand, the lower the heating temperature, the more it can reduce the amount of hydrogen in the steel, so it is preferred, but if it is too low, the finishing temperature decreases and rolling becomes difficult. Therefore, the heating temperature is preferably 950°C or more. The heating temperature is more preferably 1000°C or more. There is no particular provision for the heating time, but if it is too long, the risk of increased hydrogen introduced into the steel pipe increases, so it is preferably 180 minutes or less. The heating time is more preferably 150 minutes or less, and further preferably 120 minutes or less. There is no particular limitation on the lower limit, but the heating time is preferably 30 minutes or more, more preferably 60 minutes or more.

[0146] Rolling process

[0147] The steel billet heated in the heating step is rolled under the following conditions to form a steel pipe: The rolling can be performed by hot rolling including piercing rolling using a conventional Mannesman plug mill method or a Mannesman mandrel mill method.

[0148] Rolling end temperature: above 820℃

[0149] If the rolling end temperature is less than 820°C, the rolling load becomes excessive, increasing the risk of rolling failure. Therefore, the rolling end temperature is 820°C or higher. The rolling end temperature is preferably 850°C or higher, and more preferably 900°C or higher. On the other hand, there is no particular upper limit to the rolling end temperature, but if the temperature is too high, the metal structure tends to become non-uniform, so the rolling end temperature is preferably 1200°C or lower. The rolling end temperature is more preferably 1150°C or lower, and even more preferably 1100°C or lower.

[0150] [Cooling process (accelerated cooling process)]

[0151] In the cooling step, the steel material having the above-described composition is heated to a temperature between the Ac3 point and 1000°C, either directly or after being processed into a steel pipe, and then cooled under the cooling conditions of Group A or Group B described below. The temperature is preferably maintained for 10 minutes or longer, more preferably for 15 minutes or longer, and even more preferably for 20 minutes or longer. While the upper limit is not particularly limited, the temperature is preferably maintained for 60 minutes or shorter, and more preferably for 45 minutes or shorter.

[0152] Heating temperature after steel pipe processing: Ac3 point ~ 1000℃

[0153] If the heating temperature in the cooling process is lower than the Ac3 point, ferrite will remain in the steel after cooling, and the strength of the steel pipe and the fatigue properties in hydrogen will be reduced. Therefore, the heating temperature is above the Ac3 point. The heating temperature is preferably above Ac3 point + 30°C, more preferably above Ac3 point + 50°C. However, for the component system where Ac3 point + 30°C and Ac3 point + 50°C exceed 1000°C, the above-mentioned Ac3 point + 30°C and above Ac3 point + 50°C are not applicable. However, on the other hand, if the above-mentioned heating temperature is higher than 1000°C, the austenite grains will sometimes coarsen, causing a decrease in the impact absorption energy value and toughness of the material after heat treatment. Therefore, the above-mentioned heating temperature is set to below 1000°C. The heating temperature is preferably below 950°C, more preferably below 900°C. However, for the component system where 950°C and 900°C are lower than the Ac3 point, the above-mentioned 950°C and below 900°C are not applicable.

[0154] During the cooling process, if the temperature after rolling satisfies the present heating conditions, the steel plate may be cooled directly, or it may be reheated and then cooled after rolling. Furthermore, if the steel plate is temporarily cooled by air cooling, it may be reheated to a temperature between the Ac3 point and 1000°C and then cooled under the cooling conditions of Group A or Group B described below.

[0155] In addition, in this invention, Ac3 point (° C.) is calculated by the following formula.

[0156] Ac3(℃)=910-203[C] 1 / 2 -30[Mn]+44.7[Si]+700[P]+100[Al]+31.5[Mo]-11[Cr]-15.2[Ni]-20[Cu]+104[V]

[0157] In the formula, [M] represents the content of element M (mass %).

[0158] Average cooling rate

[0159] Group A: Cooled to 50°C or below at an average cooling rate of 15°C / s or above from 800°C to 550°C at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe, and at an average cooling rate of 15°C / s or below from 550°C to 50°C.

[0160] If the average cooling rate from 800°C to 550°C at a point 1 / 4 of the wall thickness from the inner surface of the steel pipe is less than 15°C / s, the specified carbide density cannot be achieved. While the microstructure is not particularly limited, to achieve the specified fatigue properties in hydrogen, it is preferred that either bainite or martensite be present at an area ratio of 90% or greater. If the average cooling rate is less than 15°C / s, it is difficult to obtain bainite at an area ratio of 90% or greater, and depending on the composition, martensite formation may be affected. Therefore, the average cooling rate at a point 1 / 4 of the wall thickness from the inner surface of the steel pipe is 15°C / s or greater. To suppress microstructure variation, the average cooling rate is preferably 17°C / s or greater. The average cooling rate from 800°C to 550°C is more preferably 20°C / s or greater, and most preferably 22°C / s or greater. On the other hand, to suppress grain size variation, the average cooling rate is preferably 50°C / s or less, more preferably 45°C / s or less, and even more preferably 40°C / s or less.

[0161] Furthermore, by cooling to 50°C or less under the condition that the average cooling rate from 550°C to 50°C is 15°C / s or less, retained austenite can be reduced, thereby reducing the amount of hydrogen in the steel. Therefore, the average cooling rate from 550°C to 50°C is 15°C / s or less. The average cooling rate from 550°C to 50°C is more preferably 12°C / s or less, and even more preferably 10°C / s or less. While the lower limit is not particularly limited, the average cooling rate from 550°C to 50°C is preferably 1°C / s or more.

[0162] The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling may be used alone or in combination. However, water cooling or oil cooling is preferred from 800°C to 550°C, and air cooling is preferred from 550°C to 50°C.

[0163] Group B: Cooled to 50°C or below at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe with an average cooling rate of 10°C / s or above from 800°C to 300°C and an average cooling rate of 5°C / s or below from 300°C to 50°C

[0164] If the average cooling rate from 800°C to 300°C at a point 1 / 4 of the wall thickness from the inner surface of the steel pipe is less than 10°C / s, the required carbide density cannot be achieved, resulting in deterioration in fatigue properties. Furthermore, if the average cooling rate is less than 10°C / s, it is difficult to obtain martensite with an area ratio of 90% or more, and depending on the component composition, the formation of bainite may be affected. Therefore, the average cooling rate at a point 1 / 4 of the wall thickness from the inner surface of the steel pipe is 10°C / s or higher. From the perspective of suppressing microstructure variation, the average cooling rate from 800°C to 300°C is preferably 12°C / s or higher, more preferably 15°C / s or higher, and even more preferably 17°C / s or higher. While there is no particular upper limit, the average cooling rate is preferably 60°C / s or lower.

[0165] Furthermore, by cooling to 50°C or less under the condition that the average cooling rate from 300°C to 50°C is 5°C / s or less, the amount of hydrogen in the steel can be reduced. Therefore, the average cooling rate from 300°C to 50°C is 5°C / s or less. The average cooling rate from 300°C to 50°C is preferably 1°C / s or less. The lower limit is not particularly limited, but is preferably 0.1°C / s or more.

[0166] The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling may be used alone or in combination. However, water cooling or oil cooling is preferred from 800°C to 300°C, and air cooling is preferred from 300°C to 50°C.

[0167] [Reheating and Quenching Process (Preferred Conditions)]

[0168] Reheating temperature before quenching: Ac3 point ~ 1000℃

[0169] When the temperature in the center of the plate thickness is lower than the Ac3 point, a portion of the untransformed austenite remains, and therefore the desired steel structure cannot be obtained after hot rolling, quenching, and tempering described later. Therefore, the pre-quenching heating temperature during reheating is above the Ac3 point. Preferably, it exceeds the Ac3 point. It should be noted that in order to suppress the excessive coarsening of the initial austenite grain size and improve production efficiency, the pre-quenching heating temperature is preferably below 1000°C. More preferably, it is below 980°C, and even more preferably below 960°C. Most preferably, it is below 950°C. By making the reheating temperature before quenching a temperature on the low temperature side within the range above the Ac3 point, the initial austenite grain size can be refined, and the crack propagation rate in hydrogen can be reduced.

[0170] Average cooling rate during quenching: Group A or Group B below

[0171] Group A: Cooled to 50°C or below at an average cooling rate of 15°C / s or above from 800°C to 550°C at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe, and at an average cooling rate of 15°C / s or below from 550°C to 50°C.

[0172] If the average cooling rate from 800°C to 550°C at a point 1 / 4 of the wall thickness from the inner surface of the steel pipe is less than 15°C / s, the specified carbide density cannot be achieved. Furthermore, if the average cooling rate is less than 15°C / s, it is difficult to obtain bainite with an area ratio of 90% or more, and depending on the component composition, the formation of martensite may be affected. Therefore, the average cooling rate at a point 1 / 4 of the wall thickness from the inner surface of the steel pipe is 15°C / s or more. From the perspective of suppressing microstructure variation, the average cooling rate is preferably 17°C / s or more, more preferably 20°C / s or more, and even more preferably 22°C / s or more. On the other hand, to suppress grain size variation, the average cooling rate is preferably 50°C / s or less, more preferably 47°C / s or less, and even more preferably 45°C / s or less.

[0173] Furthermore, by cooling to 50°C or below at an average cooling rate of 15°C / s or below from 550°C to 50°C, retained austenite can be reduced, thereby reducing the amount of hydrogen in the steel. Therefore, the average cooling rate from 550°C to 50°C is 15°C / s or below. The average cooling rate from 550°C to 50°C is preferably 12°C / s or below, and more preferably 10°C / s or below. While the lower limit is not particularly limited, the average cooling rate from 550°C to 50°C is preferably 1°C / s or above.

[0174] The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling may be used alone or in combination. However, water cooling or oil cooling is preferred from 800°C to 550°C, and air cooling is preferred from 550°C to 50°C.

[0175] Group B: Cooled to 50°C or below at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe with an average cooling rate of 10°C / s or above from 800°C to 300°C and an average cooling rate of 5°C / s or below from 300°C to 50°C

[0176] If the average cooling rate from 800°C to 300°C at a point 1 / 4 of the wall thickness from the inner surface of the steel pipe is less than 10°C / s, the specified carbide density cannot be achieved, resulting in deterioration in fatigue properties. Furthermore, if the average cooling rate is less than 10°C / s, it is difficult to obtain martensite with an area ratio of 90% or more, and depending on the composition, the formation of bainite may be affected. Therefore, the average cooling rate at a point 1 / 4 of the wall thickness from the inner surface of the steel pipe is 10°C / s or higher. From the perspective of suppressing microstructure variation, the average cooling rate is preferably 17°C / s or higher, more preferably 20°C / s or higher, and even more preferably 25°C / s or higher. While there is no particular upper limit for the average cooling rate, if it exceeds 60°C / s, a large amount of hard structure will form on the steel plate surface, preventing the steel structure having the target structure of the present invention from being obtained, and fatigue properties in hydrogen will be reduced. Therefore, the average cooling rate is preferably 60°C / s or lower.

[0177] Furthermore, by cooling to 50°C or less under the condition that the average cooling rate from 300°C to 50°C is 5°C / s or less, the amount of hydrogen in the steel can be reduced. Therefore, the average cooling rate from 300°C to 50°C is 5°C / s or less. The average cooling rate is preferably 3°C / s or less, and more preferably 1°C / s or less. The lower limit is not particularly limited, but is preferably 0.1°C / s or more.

[0178] The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling may be used alone or in combination. However, water cooling or oil cooling is preferred from 800°C to 300°C, and air cooling is preferred from 300°C to 50°C.

[0179] Cooling stop temperature during quenching: below 50°C

[0180] If the cooling stop temperature exceeds 50°C, the desired carbide density cannot be achieved, and the aforementioned phase transformation is not complete, so the desired steel structure cannot be obtained after tempering. Therefore, quenching is performed to a temperature of 50°C or lower. The cooling stop temperature is preferably 45°C or lower, and more preferably 40°C or lower. The lower limit is not particularly limited, but the cooling stop temperature is preferably 25°C or higher.

[0181] Tempering process

[0182] Tempering temperature: 400℃~Ac1 point

[0183] By raising the temperature under the condition of an average heating rate of 0.01°C / s or more and setting the tempering temperature to 400°C or more, it is possible to reduce austenite and hydrogen in the steel, and obtain a predetermined carbide density. The tempering temperature is preferably 450°C or more, more preferably 500°C or more. On the other hand, if the temperature is raised above the Ac1 point, there is a possibility that hydrogen in the austenite and steel will increase. Therefore, the tempering temperature is below the Ac1 point. It is preferably in the range of (Ac1 point - 30)°C or less. It should be noted that the upper limit of the above-mentioned average heating rate during tempering is not particularly limited, but is preferably 1°C / s or less. If the tempering time is too long, the carbides will coarsen and have an adverse effect on hydrogen embrittlement, so it is less than 60 minutes. The tempering time is preferably 50 minutes or less. If the tempering time is too short, the austenite in the steel will not be reduced and the hydrogen content will not be reduced, so the tempering time is preferably 10 minutes or more, more preferably 20 minutes or more.

[0184] It should be noted that, in the present invention, there is no particular provision for the method of obtaining the Ac1 point (°C), and it can be obtained, for example, as follows:

[0185] Ac1=723-14Mn+22Si-14.4Ni+23.3Cr.

[0186] In the above formula, the symbol of each element represents the content (mass %) of each element in steel, and an element not contained is represented by 0.

[0187] Dehydrogenation process

[0188] The presence of hydrogen in steel accelerates fatigue crack growth, reducing fatigue life and fatigue limit stress in hydrogen. Therefore, dehydrogenation is used to release residual hydrogen after manufacturing. Dehydrogenation reduces the amount of hydrogen in steel by holding it at high temperatures for a certain period of time before use, resulting in steel plates with excellent fatigue resistance in high-pressure hydrogen environments.

[0189] The holding time R (sec) is preferably determined based on the plate thickness of the steel material and the steel pipe, the pipe thickness t (mm), and the hydrogen diffusion coefficient D (mm) in the steel at room temperature. 2 ·sec －1 ) and adopt the following formula (A).

[0190] R≥t 2 / D···(A)

[0191] The hydrogen diffusion coefficient also varies depending on the components and metal structure, 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.

[0192] The dehydrogenation treatment 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 becomes smaller at high temperature, 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) at the temperature at which the value of D of the above-mentioned formula (A) is maintained can be used for calculation. On the other hand, when the temperature of the dehydrogenation process is too high, the material strength is significantly reduced, so the dehydrogenation treatment temperature is preferably below 550°C. The dehydrogenation treatment temperature T is more preferably below 500°C. The dehydrogenation treatment 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 treatment temperature T is preferably above room temperature. The dehydrogenation treatment temperature T is more preferably above 50°C. The dehydrogenation treatment temperature T is further preferably above 100°C, and most preferably above 150°C. The so-called dehydrogenation treatment temperature T here is the temperature of the atmosphere in the dehydrogenation treatment process. Room temperature refers to 20±10°C.

[0193] 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 steel pipe to reach the temperature of the atmosphere in the dehydrogenation treatment process (dehydrogenation treatment temperature T). Therefore, even if the above-mentioned holding time R (sec) is met at the ambient temperature, if the center of the plate thickness does not reach the dehydrogenation treatment temperature T (ambient temperature), the dehydrogenation treatment may be insufficient. Therefore, it is preferred to maintain the temperature R (sec) or more after the center of the plate thickness temperature Tc 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 hydrogen content of the steel material at the surface and the center of the plate thickness. To this end, it is preferred to maintain the temperature R (sec) specified in formula (A) at the dehydrogenation treatment temperature T, and it is further preferred to maintain the above-mentioned holding time R (sec) or more after the center of the plate thickness temperature Tc reaches the target dehydrogenation treatment temperature T. In other words, at least the former can appropriately control the hydrogen content of the steel material at the surface of the steel material and steel pipe, and when the latter is implemented, the hydrogen content of the steel material from the surface to the center of the plate thickness of the steel material and steel pipe can be appropriately controlled. The plate thickness temperature is the plate thickness center temperature Tc, which can be measured using a thermocouple or the like, or can be estimated using a finite element method or the like.

[0194] Furthermore, since scale on the steel surface hinders dehydrogenation, it is preferable to remove the scale before performing a dehydrogenation treatment. The scale removal method is not limited; for example, physical cleaning using high-pressure cleaning or chemical methods using a scale removal agent may be used. The thickness of scale removed is not limited; however, if approximately 100 μm is removed, the scale removal effect is achieved.

[0195] Second embodiment

[0196] The steel material of the present invention is described in detail below. The steel material's composition, metallurgical structure, and crack growth rate are the same as those described for the steel pipe. Regarding the manufacturing method, all steps other than the rolling and cooling steps (casting, heating, reheating and quenching, tempering, and dehydrogenation treatment) are also carried out in the same manner as those described for the steel pipe. The rolling and cooling steps are carried out as follows.

[0197] Rolling process

[0198] The steel billet heated in the heating step described in the above-mentioned method for producing a steel pipe is hot-rolled using a hot rolling mill under the following conditions.

[0199] Rolling end temperature: above 820℃

[0200] If the rolling end temperature is less than 820°C, the rolling load becomes excessive, increasing the risk of rolling failure. Therefore, the rolling end temperature is 820°C or higher. The rolling end temperature is preferably 850°C or higher, and more preferably 900°C or higher. On the other hand, there is no particular upper limit to the rolling end temperature, but if the temperature is too high, the metal structure tends to become non-uniform, so the rolling end temperature is preferably 1200°C or lower. The rolling end temperature is more preferably 1150°C or lower, and even more preferably 1100°C or lower.

[0201] Cooling process (accelerated cooling process)

[0202] In the cooling step, the steel material having the above-described composition is hot-rolled, then heated to a temperature between the Ac3 point and 1000°C and held, and then cooled under the cooling conditions of Group A or Group B described below. The temperature is preferably maintained for 10 minutes or longer, more preferably for 15 minutes or longer, and even more preferably for 20 minutes or longer. While the upper limit is not particularly limited, the temperature is preferably maintained for 60 minutes or shorter, and more preferably for 45 minutes or shorter.

[0203] Heating temperature after hot rolling: Ac3 point ~ 1000℃

[0204] If the heating temperature in the cooling process is lower than the Ac3 point, ferrite will remain in the steel after cooling, and the strength and fatigue properties of the steel will be reduced. Therefore, the heating temperature is above the Ac3 point. The heating temperature is preferably above Ac3 point + 30°C, more preferably above Ac3 point + 50°C. However, for a component system where Ac3 point + 30°C or Ac3 point + 50°C exceeds 1000°C, the above-mentioned Ac3 point + 30°C or above, Ac3 point + 50°C or above are not applicable. On the other hand, if the above-mentioned heating temperature is higher than 1000°C, the austenite grains will sometimes coarsen, causing a decrease in the impact absorption energy value and toughness of the material after heat treatment. Therefore, the above-mentioned heating temperature is set to below 1000°C. It is more preferably below 950°C, and further preferably below 900°C. However, for a component system where 950°C or 900°C is lower than the Ac3 point, the above-mentioned 950°C or below, 900°C or below are not applicable.

[0205] In the cooling process here, if the temperature after rolling meets the present heating conditions, the steel may be cooled directly, or it may be heated again and then cooled after rolling. In addition, if the steel is temporarily cooled by air cooling, it may be heated again to a temperature of Ac3 point to 1000°C, and then cooled under the cooling conditions of the following Group A or Group B (this case is referred to as quenching).

[0206] In addition, in this invention, Ac3 point (° C.) is calculated by the following formula.

[0207] Ac3(℃)=910-203[C] 1 / 2 -30[Mn]+44.7[Si]+700[P]+100[Al]+31.5[Mo]-11[Cr]-15.2[Ni]-20[Cu]+104[V]

[0208] In the formula, [M] represents the content of element M (mass %).

[0209] Average cooling rate

[0210] Group A: Cooled to 50°C or below at an average cooling rate of 15°C / s or above from 800°C to 550°C at a position 1 / 4 of the thickness from the steel surface, and at an average cooling rate of 15°C / s or below from 550°C to 50°C.

[0211] When the average cooling rate from 800°C to 550°C at a position of 1 / 4 of the thickness of the plate from the surface of the steel material is less than 15°C / s, the specified carbide density cannot be obtained. In addition, the structure is not particularly limited, but in order to obtain the specified fatigue properties in hydrogen, it is preferred to obtain either bainite or martensite with an area ratio of 90% or more. When the above-mentioned average cooling rate is less than 15°C / s, it is difficult to obtain bainite with an area ratio of 90% or more, but the formation of martensite is sometimes affected depending on the component composition. Therefore, the average cooling rate at a position of 1 / 4 of the thickness of the plate from the surface of the steel material is 15°C / s or more. From the viewpoint of suppressing the deviation of the structure, the average cooling rate is preferably 17°C / s or more. It is more preferably 20°C / s or more, and further preferably 22°C / s or more. On the other hand, in order to suppress the deviation of the grain size, the average cooling rate is 50°C / s or less. It is preferably 47°C / s or less, and more preferably 45°C / s or less.

[0212] Furthermore, by cooling to 50°C or less at an average cooling rate of 15°C / s or less from 550°C to 50°C, retained austenite can be reduced, thereby reducing the amount of hydrogen in the steel. Therefore, the average cooling rate from 550°C to 50°C is 15°C / s or less. While the lower limit is not particularly limited, the average cooling rate from 550°C to 50°C is preferably 1°C / s or greater.

[0213] The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling may be used alone or in combination. However, water cooling or oil cooling is preferred from 800°C to 550°C, and air cooling is preferred from 550°C to 50°C.

[0214] Group B: Cooled to 50°C or below at a position 1 / 4 of the thickness from the steel surface with an average cooling rate of 10°C / s or above from 800°C to 300°C and an average cooling rate of 5°C / s or below from 300°C to 50°C

[0215] If the average cooling rate from 800°C to 300°C at a position 1 / 4 of the thickness of the plate from the surface of the steel material is less than 10°C / s, the specified carbide density cannot be obtained, resulting in deterioration of fatigue properties. In addition, if the above-mentioned average cooling rate is less than 10°C / s, it is difficult to obtain martensite with an area ratio of more than 90%, and depending on the component composition, it sometimes affects the formation of bainite. Therefore, the average cooling rate at a position 1 / 4 of the thickness of the plate from the surface of the steel material is 10°C / s or more. From the viewpoint of suppressing the deviation of the structure, it is more preferably 12°C / s or more. The average cooling rate is more preferably 15°C / s or more, and further preferably 17°C / s or more. On the other hand, the upper limit of the above-mentioned average cooling rate is not particularly specified, but if it exceeds 60°C / s, a large amount of hard structure will be generated on the surface of the steel plate, and the steel structure with the target structure in the present invention cannot be obtained, and the fatigue properties in hydrogen are reduced. Therefore, the above-mentioned average cooling rate is preferably 60°C / s or less.

[0216] Furthermore, by cooling the steel to 50°C or less under the condition of an average cooling rate of 5°C / s or less from 300°C to 50°C, the amount of hydrogen in the steel can be reduced. Therefore, the average cooling rate from 300°C to 50°C is 5°C / s or less. The average cooling rate is preferably 1°C / s or less, and more preferably 0.8°C / s or less. The lower limit is not particularly limited, but is preferably 0.1°C / s or more.

[0217] The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling may be used alone or in combination. However, water cooling or oil cooling is preferred from 800°C to 300°C, and air cooling is preferred from 300°C to 50°C.

[0218] This is not necessary in the case of a thick plate, but it is preferable to wind a thin steel plate into a coil.

[0219] It should be noted that the steel with excellent fatigue properties in hydrogen of the present invention is a steel having the above-mentioned composition, and includes various categories such as thin plates, thick plates, steel pipes, etc. with excellent fatigue crack growth resistance in hydrogen, or can be a steel for hydrogen pipelines formed into a predetermined shape.

[0220] Under the above conditions, a steel pipe and a steel material having excellent fatigue properties in hydrogen that satisfy a predetermined crack growth rate in hydrogen can be obtained.

[0221] Example 1

[0222] The following examples demonstrate the effects of the present invention. While the following examples illustrate preferred embodiments of the present invention, the present invention is not limited in any way by the following examples. It should be noted that the following examples were conducted using the manufacturing method and property evaluation of seamless steel pipes used in actual steel structures.

[0223] Billets with the compositions shown in Tables 1-1 and 1-2, steel pipes Nos. 1 to 29 (billets A to AC) and 40 to 87 (billets AN to CI), were produced at a casting speed of 0.6 m / min. These billets were heated to 1250°C and expanded to produce seamless steel pipes. These steel pipes were produced under conditions where expansion was completed at or above 820°C. The resulting steel pipes were heated and held at 950°C for pipes with an Ac3 point of 950°C or less, and at 1000°C for pipes with an Ac3 point exceeding 950°C. They were then water-cooled under the conditions shown in Tables 2-1 and 2-2 and tempered. The metallographic structure and mechanical properties were evaluated. Furthermore, slabs with the compositions shown in Tables 1-1 and 1-2, steel material Nos. 30 to 39 (slabs AD to AM), were produced at a casting speed of 0.6 m / min, heated to 1250°C, and then rolled using a hot rolling mill at or above 820°C. The resulting steels were heated and held at 950°C for steels with an Ac3 point of 950°C or less, and at 1000°C for steels with an Ac3 point exceeding 950°C. The resulting steels were then water-cooled under the conditions listed in Tables 2-1 and 2-2, and then tempered. The metallurgical structure and mechanical properties were evaluated in the same manner as for the steel pipes. After tempering, Nos. 2, 5, 14, 15, 43, and 63 to 69 were also subjected to a dehydrogenation treatment. The dehydrogenation treatment was performed by holding the dehydrogenation treatment temperature T at 50°C as the ambient temperature for 3 hours, followed by natural cooling. The evaluation method is as follows. The tempering temperature was arbitrarily adjusted so that the tensile strength of the material was within the range of 520 to 700 MPa.

[0224] Furthermore, billets with the compositions shown in Table 2-3 for Steel Pipes Nos. 88 to 101 (Bills Nos. AO1 to BB1) were produced at various casting speeds. These billets were heated to 1250°C and expanded to produce seamless steel pipes. The compositions of Billets Nos. AO1 to BB1 were identical to those of Billets Nos. AO to BB shown in Table 1-2. These steel pipes were produced under conditions where expansion was completed at or above 820°C. The resulting steel pipes were heated and held at 950°C for pipes with an Ac3 point of 950°C or less, and at 1000°C for pipes with an Ac3 point exceeding 950°C. The pipes were then water-cooled under the conditions shown in Table 2-3 and tempered under the conditions shown in Table 2-3. The metallographic structure and mechanical properties were evaluated using the following evaluation method. The tempering temperature was arbitrarily adjusted to achieve a tensile strength within the range of 520 MPa to 700 MPa. The dehydrogenation treatment was performed by holding the dehydrogenation treatment temperature (T) at 50°C for 3 hours, followed by natural cooling.

[0225] The fatigue crack propagation characteristics were investigated and evaluated by fatigue crack growth tests. From each steel material, a CT (compact tensile) 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, and a fatigue test was performed at a frequency of 1 Hz, a repetitive waveform of a sine wave, and a stress ratio of R = 0.1 to obtain the fatigue crack propagation characteristics. The fatigue crack length was measured by the compliance method using a clamp gauge, and the fatigue crack propagation velocity in 5 MPa high-pressure hydrogen was obtained. The test was performed at room temperature (20 ± 10 ° C). It should be noted that for the test piece, when the plate thickness is less than 10 mm, it is ground from the surface every 0.5 mm to make 2 mm, 5 mm, 8 mm, and 9 mm respectively. In the case of plate thicknesses other than these, a 10 mm thick test piece was taken from the position of t / 2 (t: plate thickness), and the front and back of the crack extension part were mirror polished. 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) within the test specimen was evaluated as a representative value. The results are shown in Tables 2-1, 2-2, and 2-3.

[0226] The method for measuring carbides in steel is as follows. A test piece was cut from the center of the steel plate thickness, parallel to the plate thickness direction, and etched with nitric alcohol. Carbides were observed using an SEM. Ten fields of view were randomly selected at an accelerating voltage of 15 kV and a magnification of 20,000 times, and observed. The average value of the ten fields of view was used as the number of carbides. Carbides with a diameter of 200 nm or more were counted as 20 per 10 μm. 2 The following cases are Y, more than 20 / 10μm 2 The case is N, as shown in Tables 2-1, 2-2, and 2-3.

[0227] The method for measuring the austenite content of steel materials is as follows.

[0228] Samples for metallographic observation were collected from the center of the longitudinal width of the steel materials and steel pipes obtained in accordance with the above-described procedures. A cross section parallel to the longitudinal direction was polished as the observation surface. The surface layer was then removed by chemical polishing using picric acid etching. X-ray diffraction analysis was then performed. Specifically, a Co-Kα source was used as the incident X-ray source, and the area fraction of retained austenite was calculated based on the intensity ratio of the (200), (211), and (220) planes of ferrite to the (200), (220), and (311) planes of austenite.

[0229] Hydrogen heating analysis

[0230] The amount of hydrogen remaining in the steel was determined using a temperature desorption analysis method using a low-temperature temperature-type hydrogen analyzer (gas chromatograph type) (JTF-20AL). The temperature desorption analysis was performed at a heating rate of 200°C / h over a temperature range from room temperature to 400°C, and the total amount was used as the hydrogen content. The test specimen was a cylindrical shape with a length of 30 mm and a diameter of 7 Φ along the longitudinal direction of the steel pipe, measured at a position 1 / 4 of the thickness of the steel plate and 1 / 4 of the distance from the inner surface of the steel pipe. It should be noted that this hydrogen content is the hydrogen content before the high-pressure hydrogen fatigue test described below, and is the H content shown in Tables 1-1 and 1-2.

[0231] In addition, the steel pipe was also subjected to the above-mentioned various tests by the same method as the above-mentioned steel material.

[0232] All the examples of the present invention meet the fatigue crack growth rate in hydrogen of 1.0×10 －6 m / cycle or less.

[0233] Note that Steel Pipes No. 94 and 101 in Table 2-3, whose casting speeds were outside the range of the present invention, had fatigue crack growth rates in hydrogen outside the range of the present invention due to the generation of coarse inclusions, and thus serve as comparative examples.

[0234] [Table 1-1]

[0235]

[0236] [Table 1-2]

[0237]

[0238] [Table 2-1]

[0239]

[0240] Underline: indicates outside the scope of the present invention. γ: Austenite

[0241] [Table 2-2]

[0242]

[0243] Underline: indicates outside the scope of the present invention. γ: Austenite

[0244] [Table 2-3]

[0245]

[0246] Underline: indicates outside the scope of the present invention. γ: Austenite

[0247] Example 2

[0248] The following describes examples that verify the effects of the present invention. It should be noted that in the following examples, steel pipes were manufactured under the following manufacturing conditions and their properties were evaluated. Steel pipes having the same composition as billets No. Q and BC shown in Tables 1-1 and 1-2, and AS1 shown in Table 2-3 were used, and the cooling process was carried out under the prescribed conditions. After the cooling process (before the tempering process), the steel pipes were reheated under the conditions of Table 3, and the quenching process was performed. The properties of the obtained steel pipes were evaluated. Steel pipes No. 17A to 17C shown in Table 3 were obtained by reheating steel pipe No. 17 shown in Tables 1-1 and 2-1. In addition, steel pipes No. 55A to 55C were obtained by reheating steel pipe No. 55 shown in Tables 1-2 and 2-2, and steel pipes No. 92A to 92B were obtained by reheating steel pipe No. 92 shown in Table 2-3.

[0249] All the examples in Example 2 satisfy the crack growth rate da / dN in hydrogen of 1.0×10 －6 Among them, when the reheating and quenching steps are carried out under more preferable conditions, the crack propagation characteristics are more excellent.

[0250] [Table 3]

[0251]

[0252] Example 3

[0253] The following describes examples that demonstrate the effects of the present invention. In the following examples, steel pipes were produced under the following manufacturing conditions, and their properties were evaluated. Using billets No. N and AQ shown in Tables 1-1 and 1-2, and AX1 shown in Table 2-3, the steel pipes were produced under the same conditions as for Steel Pipes Nos. 14 and 43 shown in Tables 2-1 and 2-2, and Steel Pipe No. 97 shown in Table 2-3, up to the tempering step. Properties were evaluated when the dehydrogenation treatment conditions were varied. The results are shown in Table 4.

[0254] In Example 1, the dehydrogenation treatment of Steel Pipes Nos. 14, 43, and 97 was carried out at a dehydrogenation treatment temperature T (atmosphere temperature) of 50°C and a holding time of 3 hours. However, in this Example, Steel Pipes Nos. 14D, 43D, and 97D were treated so that the dehydrogenation treatment temperature T (atmosphere temperature) was 50°C and the holding time tc after the plate center temperature Tc reached 50°C satisfied Formula (A). Steel Pipes Nos. 14E, 43E, and 97E were treated so that the dehydrogenation treatment temperature T (atmosphere temperature) was 50°C and the holding time tc after the plate center temperature Tc reached 50°C satisfied Formula (A). However, the holding time tc after the plate center temperature Tc reached 50°C did not satisfy Formula (A).

[0255] For steel pipes No. 14F and 97F, 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).

[0256] In Table 4, "Dehydrogenation holding time t is Y" means that the dehydrogenation treatment temperature T (ambient temperature) is 50°C and the holding time t satisfies equation (A). "Dehydrogenation holding time t is N" means that the dehydrogenation treatment temperature T (ambient temperature) is 50°C but the holding time t does not satisfy equation (A). Furthermore, "Steel material center temperature Tc holding time tc is Y" means that the holding time tc after the plate center temperature Tc reaches 50°C satisfies equation (A), while "Steel material center temperature Tc holding time tc is N" means that the plate center temperature Tc reaches 50°C but the holding time tc after Tc reaches 50°C does not satisfy equation (A).

[0257] The fatigue crack propagation characteristics were evaluated by the fatigue crack growth test described in Example 1.

[0258] All the examples of the present invention satisfy the crack growth rate da / dN in hydrogen of 1.0×10 －6 Among them, the more preferable the dehydrogenation treatment conditions are, the better the crack propagation characteristics are.

[0259] [Table 4]

[0260]

[0261] γ: Austenite

Claims

1. A steel pipe having excellent fatigue properties in hydrogen, comprising, by mass%, 0.10-0.45% C, 0.01-2.0% Si, 0.3-2.0% Mn, 0.01-0.15% Al, 0.0005-0.008% N, 0.015% or less P, 0.0015% or less S, 0.01% or less O, 0.01% or less H, 0.0010% or less Cu, %, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, Nb: 0-0.5%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.3%, 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; The retained austenite is less than 3%, The number of carbides with a diameter of 200 nm or more is 20 / 10 μm 2 the following, The stress expansion coefficient range in hydrogen above 1MPa = crack growth rate da / dN at 20MPa√m is 1.0×10 － 6 m·cycle -1 the following.

2. A method for manufacturing a steel pipe, comprising: A casting step of casting the steel billet having the component composition of claim 1 at a casting speed of 1.8 m / min or less, The heating process is carried out at a temperature below 1350°C. In the hot rolling step, the steel billet heated in the heating step is rolled at a rolling end temperature of 820° C. or higher to form a steel pipe. A cooling step, after the steel pipe obtained in the hot rolling step is maintained at a temperature of Ac3 point to 1000°C, the cooling conditions are the following group A or group B, and A tempering step, wherein the steel pipe obtained in the cooling step is tempered at 400°C to Ac1 point for less than 60 minutes; Group A: The steel pipe is cooled to 50°C or below at an average cooling rate of 15°C / s or above at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe from 800°C to 550°C, and at an average cooling rate of 15°C / s or below at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe from 550°C to 50°C, Group B: The steel pipe is cooled to 50°C or less at an average cooling rate from 800°C to 300°C of 10°C / s or more at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe, and at an average cooling rate from 300°C to 50°C of 5°C / s or less at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe.

3. The method for manufacturing a steel pipe according to claim 2, wherein: There is a quenching process before the tempering process, which reheats the steel pipe to Ac3 point ~ 1000°C, and the cooling conditions are the following group A or group B; Group A: The steel pipe is cooled to 50°C or below at an average cooling rate of 15°C / s or above at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe from 800°C to 550°C, and at an average cooling rate of 15°C / s or below at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe from 550°C to 50°C, Group B: The steel pipe is cooled to 50°C or less at an average cooling rate from 800°C to 300°C of 10°C / s or more at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe, and at an average cooling rate from 300°C to 50°C of 5°C / s or less at a position 1 / 4 of the wall thickness from the inner surface of the steel pipe.

4. The method for manufacturing a steel pipe according to claim 2 or 3, wherein: The casting speed is below 1.0 m / min.

5. A steel material having excellent fatigue properties in hydrogen, comprising, by mass%, C: 0.10-0.45%, Si: 0.01-2.0%, Mn: 0.3-2.0%, Al: 0.01-0.15%, N: 0.0005-0.008%, P: 0.015% or less, S: 0.0015% or less, O: 0.01% or less, H: 0.0010% or less, Cu: 0-2.5% %, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, Nb: 0-0.5%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.3%, 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; The retained austenite is less than 3%, The number of carbides with a diameter of 200 nm or more is 20 / 10 μm 2 the following, The stress expansion coefficient range in hydrogen above 1MPa = crack growth rate da / dN at 20MPa√m is 1.0×10 － 6 m·cycle -1 the following.

6. A method for manufacturing a steel material, comprising: A casting step of casting the steel billet having the component composition according to claim 5 at a casting speed of 1.8 m / min or less, The heating process is carried out at a temperature below 1350°C. The hot rolling step is to roll the steel billet heated in the heating step at a rolling end temperature of 820° C. or higher. A cooling step, after the steel material obtained in the hot rolling step is maintained at a temperature of Ac3 point to 1000°C, the cooling conditions are the following group A or group B, and A tempering step, wherein the steel obtained in the cooling step is tempered at 400°C to Ac1 point for less than 60 minutes; Group A: The steel is cooled to 50°C or below at an average cooling rate of 15°C / s or above at a position 1 / 4 of the wall thickness from the surface of the steel material, and at an average cooling rate of 15°C / s or below at a position 1 / 4 of the wall thickness from the surface of the steel material, from 550°C to 50°C. Group B: The steel is cooled to below 50°C with an average cooling rate from 800°C to 300°C of 10°C / s or more at a position 1 / 4 of the wall thickness from the steel surface, and an average cooling rate from 300°C to 50°C of 5°C or less at a position 1 / 4 of the wall thickness from the steel surface.

7. The method for manufacturing a steel material according to claim 6, wherein: There is a quenching process before the tempering process, which reheats the steel material to Ac3 point ~ 1000°C, and the cooling conditions are the following group A or group B; Group A: The steel is cooled to 50°C or below at an average cooling rate of 15°C / s or above at a position 1 / 4 of the wall thickness from the surface of the steel material, and at an average cooling rate of 15°C / s or below at a position 1 / 4 of the wall thickness from the surface of the steel material, from 550°C to 50°C. Group B: The steel is cooled to below 50°C with an average cooling rate from 800°C to 300°C of 10°C / s or more at a position 1 / 4 of the wall thickness from the steel surface, and an average cooling rate from 300°C to 50°C of 5°C or less at a position 1 / 4 of the wall thickness from the steel surface.

8. The method for producing a steel material according to claim 6 or 7, wherein: The casting speed is below 1.0 m / min.

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