WO2024071358A1 - High-strength line pipe steel material having excellent fracture toughness in hydrogen, method for manufacturing same, steel tube for high-strength line pipes, and method for manufacturing same - Google Patents

High-strength line pipe steel material having excellent fracture toughness in hydrogen, method for manufacturing same, steel tube for high-strength line pipes, and method for manufacturing same Download PDF

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WO2024071358A1
WO2024071358A1 PCT/JP2023/035560 JP2023035560W WO2024071358A1 WO 2024071358 A1 WO2024071358 A1 WO 2024071358A1 JP 2023035560 W JP2023035560 W JP 2023035560W WO 2024071358 A1 WO2024071358 A1 WO 2024071358A1
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hydrogen
temperature
steel
content
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PCT/JP2023/035560
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French (fr)
Japanese (ja)
Inventor
佳宏 西原
拓史 岡野
奈穂 井上
大地 泉
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Ｊｆｅスチール株式会社
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Publication of WO2024071358A1 publication Critical patent/WO2024071358A1/en

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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/10—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur

Definitions

  • the present invention relates to a high-strength line pipe steel material that has excellent fracture toughness in hydrogen in a high-pressure hydrogen gas environment of 1 MPa or more, and is suitable for applications such as line pipes for transporting hydrogen gas, a manufacturing method thereof, and a high-strength line pipe steel pipe and a manufacturing method thereof.
  • Austenitic stainless steels such as SUS316L, which exhibit fracture toughness in hydrogen, are used for steel structures used in high-pressure hydrogen gas environments.
  • the steel has low strength, and when designed to withstand high hydrogen pressure, the wall thickness becomes thicker and the line pipe becomes very expensive, making it unsuitable for pipeline construction. For this reason, there has been a demand for steel materials for hydrogen line pipes that are lower cost and can withstand high-pressure hydrogen gas environments.
  • Patent Document 3 proposes an austenitic steel material with a high Mn content.
  • the technology described in Patent Document 3 makes it possible to provide a steel material that is less expensive than austenitic stainless steel, but because it is austenitic, it is more expensive than low-alloy steel.
  • no consideration is given to suppressing pitting corrosion, which is the starting point of hydrogen-induced cracking, such as HIC resistance and SSCC resistance.
  • the limit of fatigue failure of a steel structure used in a high-pressure hydrogen gas environment corresponds to the critical crack length calculated from the operating conditions of the pipeline and the hydrogen-induced crack propagation lower limit K IH , which corresponds to the fracture toughness value of the steel in hydrogen gas. From the viewpoint of extending the life and improving the safety of hydrogen structures, increasing the K IH of the steel is considered to be one effective guideline.
  • Patent Document 4 proposes a manufacturing method for steel materials with excellent K IH , but does not mention the characteristics of the welds. In general, welds are more susceptible to property deterioration due to hydrogen than base materials. Therefore, it is important to improve K IH , including welds.
  • the present invention has been made to solve the above-mentioned problems of the prior art, and aims to provide a high-strength line pipe steel material with excellent fracture toughness in hydrogen in a high-pressure hydrogen gas environment, suitable for steel structures used in a high-pressure hydrogen gas environment, such as line pipes for 100% hydrogen gas or natural gas (natural gas is a gas whose main components are hydrocarbons such as methane and ethane) containing hydrogen at a hydrogen partial pressure of 1 MPa or more, a manufacturing method thereof, and a high-strength line pipe steel pipe and a manufacturing method thereof.
  • the high-pressure hydrogen gas environment is assumed to be high-pressure hydrogen gas of 1 MPa or more, or an environment containing 0.2% or more hydrogen gas.
  • the term "excellent fracture toughness in hydrogen under high pressure hydrogen gas environment” refers to a case where the hydrogen-induced crack propagation lower limit K IH is 80 MPa ⁇ m 1/2 or more, which is determined by conducting a fracture toughness test under both environments of room temperature (20 ⁇ 10 °C), hydrogen gas at a pressure of 1 MPa or more, or a mixed atmosphere of natural gas (mainly composed of hydrocarbons such as methane and ethane) containing hydrogen at a hydrogen partial pressure of 1 MPa or more.
  • the fracture toughness value refers to a value determined by conducting a fracture toughness test in accordance with ASTM E399, ASTM E1820, and ASTM E1681.
  • 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.
  • steel here includes thin steel plates, thick steel plates, seamless steel pipes, electric resistance welded steel pipes, steel sections, steel bars, etc.
  • the present inventors conducted technical studies on the conditions that should be satisfied by a steel material for obtaining a high-strength linepipe steel material and a high-strength linepipe steel pipe having excellent fracture toughness in hydrogen under a high-pressure hydrogen gas environment, with the aim of suppressing hydrogen absorption into the steel material, which is the root cause of hydrogen embrittlement.
  • the hydrogen-induced crack propagation lower limit K IH of the steel material and the steel pipe is improved in a metal structure in which the number of inclusions having an aspect ratio of 2.0 or more and a length of 10 ⁇ m or more is 15 pieces/100 mm2 or less, and the maximum grain size of the bainite in the range from the surface of the steel material and the steel pipe to the center of the plate thickness is 25 ⁇ m or less.
  • the hydrogen-induced crack propagation lower limit K IH of the steel material is further improved if the area fraction of the retained austenite is 0 to 3%, and the area fraction of the bainite in the range from the surface of the steel material and the steel pipe to the center of the plate thickness is 90% or more.
  • high strength refers to a tensile strength of 520 MPa or more.
  • the gist of the present invention is as follows. [1] In mass%, C: 0.02 to 0.15%, Si: 0.01 to 2.0%, Mn: 0.5 to 1.5%, P: 0.0001 to 0.015%, S: 0.0002 to 0.0015%, Al: 0.005 to 0.15%, O: 0.01% or less, N: 0.010% or less, Nb: 0.10% or less, H: 0.02 ppm or less, Or even more so: Ca: 0 to 0.005%, Ni: 0 to 2.0%, Ti: 0 to 0.1%, Cu: 0 to 1.0%, Cr: 0 to 1.0%, Mo: 0 to 0.60%, W: 0 to 1.0%, V: 0 to 0.10%, Zr: 0 to 0.050%, Mg: 0 to 0.01%, REM: 0 to 0.01%, B: 0 to 0.0020%, Ta: 0 to 0.2%, Hf: 0 to 0.2%, Re: 0 to 0.005%, Sn: 0
  • the chemical composition comprises, in mass%, Ca: 0.0001 to 0.005%, Ni: 0.01 to 2.0%, Ti: 0.005 to 0.1%, Cu: 0.01 to 1.0%, Cr: 0.01 to 1.0%, Mo: 0.01 to 0.60%, W: 0.01 to 1.0%, V: 0.01 to 0.10%, Zr: 0.0001 to 0.050%, Mg: 0.0001 to 0.01%, REM: 0.0001 to 0.01%, B: 0.0001 to 0.0020%, Ta: 0.0001 to 0.2%, Hf: 0.0001 to 0.2%, Re: 0.0001 to 0.005%, Sn: 0.0001 to 0.3%,
  • a steel material for high-strength linepipes having excellent fracture toughness in hydrogen according to [1] or [2], wherein the area fraction of retained austenite is 0 to 3%, and the area fraction of the bainite in the range from the steel material surface to the center of the plate thickness is 90% or more.
  • the present invention relates to a method for producing a high-strength steel material for line pipes having excellent fracture toughness in hydrogen.
  • the chemical composition comprises, in mass%, Ca: 0.0001 to 0.005%, Ni: 0.01 to 2.0%, Ti: 0.005 to 0.1%, Cu: 0.01 to 1.0%, Cr: 0.01 to 1.0%, Mo: 0.01 to 0.60%, W: 0.01 to 1.0%, V: 0.01 to 0.10%, Zr: 0.0001 to 0.050%, Mg: 0.0001 to 0.01%, REM: 0.0001 to 0.01%, B: 0.0001 to 0.0020%, Ta: 0.0001 to 0.2%, Hf: 0.0001 to 0.2%, Re: 0.0001 to 0.005%, Sn: 0.0001 to 0.3%,
  • steel materials with extremely improved fracture toughness in hydrogen under high-pressure hydrogen gas environments can be easily and simply manufactured, which is of great industrial benefit.
  • the hydrogen absorption resistance characteristics of steel structures such as high-pressure hydrogen gas line pipes can be significantly improved, which also has the effect of greatly contributing to improving the safety of steel structures.
  • a steel material will be specifically described, then as a second embodiment, a UOE steel pipe, which is an example of a steel pipe of the present invention, will be specifically described, and as a third embodiment, an electric resistance welded steel pipe, which is an example of a steel pipe of the present invention, will be specifically described.
  • C 0.02 to 0.15%
  • the C content is set to 0.02% or more.
  • the C content is 0.03% or more. More preferably, the C content is 0.035% or more. Even more preferably, the C content is 0.04% or more.
  • the C content exceeds 0.15%, the weldability decreases. For this reason, the C content is limited to 0.15% or less.
  • the C content is 0.10% or less.
  • the C content exceeds 0.08%, the hardness of the surface layer and the central segregation increases during controlled cooling, so that the SSCC resistance and HIC resistance may deteriorate. Furthermore, the toughness also deteriorates. For this reason, the C content is more preferably 0.08% or less. Even more preferably, the C content is 0.06% or less.
  • Si 0.01 to 2.0% Si is contained for deoxidation, but if the content is less than 0.01%, the deoxidation effect is insufficient, so the Si content is set to 0.01% or more.
  • the Si content is preferably 0.02% or more. More preferably, the Si content is 0.05% or more. Even more preferably, the Si content is 0.08% or more. Since the above effect is observed up to 2.0%, the Si content is set to 2.0% or less.
  • the Si content is preferably 1.8% or less, more preferably 1.5% or less. Even more preferably, the Si content is 1.0% or less. However, if the Si content exceeds 0.5%, toughness and weldability may be deteriorated, so the Si content is most preferably 0.5% or less.
  • Mn 0.5 to 1.5% Mn effectively contributes to improving strength and toughness, but if the content is less than 0.5%, the effect of inclusion is poor, so the Mn content is set to 0.5% or more.
  • the Mn content is 0.6% or more, more preferably 0.8% or more. More preferably, the Mn content is 1.0% or more.
  • the Mn content is limited to 1.5% or less.
  • the Mn content is 1.4% or less. More preferably, the Mn content is 1.3% or less, and even more preferably, the Mn content is 1.2% or less.
  • P 0.0001 to 0.015%
  • P is an inevitable impurity element that deteriorates weldability and increases the hardness of the central segregation, thereby deteriorating HIC resistance. Since this tendency becomes significant when the P content exceeds 0.015%, the P content is limited to 0.015% or less.
  • the P content is preferably 0.012% or less, and more preferably 0.010% or less. More preferably, the P content is 0.008% or less. The lower the content, the better, but from the viewpoint of refining costs, the P content is set to 0.0001% or more.
  • S 0.0002 to 0.0015%
  • S is an inevitable impurity element, and since it becomes MnS inclusions in steel and deteriorates HIC resistance, it is preferable that the S content is small, but up to 0.0015% is permissible. Therefore, the S content is set to 0.0015% or less.
  • the S content is preferably 0.0010% or less, and more preferably 0.0008% or less. The lower the content, the better, but from the viewpoint of refining costs, it is set to 0.0002% or more.
  • Al 0.005 to 0.15%
  • Al is added as a deoxidizer, but if it is less than 0.005%, there is no effect, so the Al content is 0.005% or more. On the other hand, if it exceeds 0.15%, the cleanliness of the steel decreases and the toughness deteriorates, so the Al content is set to 0.15% or less.
  • the Al content is preferably 0.12% or less, more preferably 0.10% or less. More preferably, the Al content is 0.08% or less.
  • O 0.01% or less O is a cause of oxide inclusions, so the less the better. This effect does not become a problem if the O content is 0.01% or less, so the O content is set to 0.01% or less.
  • the O content is preferably 0.0080% or less. More preferably, the O content is less than 0.0030%.
  • the lower limit is not particularly limited, but may be 0.0005% or more.
  • N 0.010% or less N effectively contributes to improving strength, but if the content exceeds 0.010%, the hardness increases during controlled cooling, resulting in deterioration of toughness. For this reason, the N content is set to 0.010% or less.
  • the N content is preferably set to 0.008% or less, more preferably set to 0.006% or less, and even more preferably set to 0.004% or less.
  • the content is preferably set to 0.00001% or more. More preferably, the N content is 0.002% or more.
  • Nb 0.10% or less
  • Nb is an element effective for increasing the strength and toughness of steel. If the content is less than 0.001%, the effect of the content is poor, so 0.001% or more is preferable. On the other hand, if the content exceeds 0.10%, the toughness of the welded part deteriorates, so the Nb content is set to 0.10% or less.
  • the Nb content is preferably set to 0.095% or less.
  • the Nb content is more preferably set to 0.090% or less, and even more preferably set to 0.085% or less.
  • the Nb content is most preferably set to 0.080% or less.
  • H 0.02 ppm or less H may be introduced into the steel material in various processes during manufacturing, and if the amount introduced is large, the risk of cracking after solidification increases and the K IH may be significantly reduced. These effects are not a problem if the amount is 0.02 ppm or less, so the H content is set to 0.02 ppm or less.
  • the H content is preferably 0.015 ppm or less, more preferably 0.008 ppm or less.
  • the H content is further preferably 0.005 ppm or less, and most preferably less than 0.002 ppm.
  • the lower limit is not particularly limited, but is preferably 0.0008 ppm or more for reasons of manufacturing costs.
  • the H content is more preferably 0.001 ppm or more.
  • the amount of hydrogen is the amount of hydrogen remaining after forming of steel material, steel pipe, UOE, etc.
  • the chemical composition disclosed herein may further contain one or more elements selected from Ca, Ni, Ti, Cu, Cr, Mo, W, V, Zr, Mg, REM, B, Ta, Hf, Re, Sn, and Sb in the following ranges:
  • Ca 0 to 0.005% Since Ca is an element effective in improving HIC resistance by controlling the morphology of sulfide-based inclusions, when Ca is contained, the Ca content may be 0% or more, but if it is less than 0.0001%, the effect of adding it is insufficient. Therefore, when Ca is contained, the Ca content is preferably 0.0001% or more. More preferably, it is 0.0005% or more. On the other hand, if it exceeds 0.005%, not only the effect is saturated, but also the HIC resistance is deteriorated due to the decrease in the cleanliness of the steel, so when Ca is contained, the Ca content is limited to 0.005% or less. The Ca content is preferably 0.004% or less. The Ca content is more preferably 0.002% or less, and even more preferably 0.0008% or less.
  • Ni 0 to 2.0%
  • Ni is an element effective in improving toughness and increasing strength, and when Ni is contained, the Ni content may be 0% or more, but to obtain this effect, it is preferable to contain 0.01% or more.
  • the Ni content is more preferably 0.1% or more.
  • the Ni content is 2.0% or less.
  • the Ni content is preferably 1.8% or less.
  • the Ni content is more preferably 1.4% or less, and even more preferably 0.8% or less.
  • Ti 0 to 0.1% Since Ti contributes to increasing the strength of the steel material, when Ti is contained, the Ti content may be 0% or more. In order to obtain the above effect, when Ti is contained, the content is preferably 0.005% or more. More preferably, it is 0.008% or more. On the other hand, when the content exceeds 0.1%, the effect is saturated and becomes a factor of increasing costs, so when Ti is contained, the Ti content is 0.1% or less.
  • the Ti content is preferably 0.08% or less, and more preferably 0.06% or less. In order to suppress costs, the Ti content is further preferably 0.05% or less. The Ti content is most preferably 0.04% or less.
  • Cu 0 to 1.0%
  • Cu is an element effective in improving toughness and increasing strength, and when Cu is contained, the Cu content may be 0% or more, but to obtain this effect, it is preferable to contain 0.01% or more. It is more preferable to make it 0.05% or more.
  • the Cu content is 1.0% or less.
  • the Cu content is preferably 0.95% or less, and more preferably 0.9% or less. More preferably, the Cu content is 0.85% or less. Most preferably, the Cu content is 0.5% or less.
  • Cr 0 to 1.0% Like Mn, Cr is an effective element for obtaining sufficient strength even with low C.
  • the Cr content may be 0% or more, but to obtain this effect, it is preferable to contain 0.01% or more. It is more preferable to make it 0.05% or more.
  • the content is too high, the hardenability becomes excessive, so that the SSCC resistance deteriorates. In addition, the weldability also deteriorates. Therefore, when Cr is contained, it is 1.0% or less.
  • the Cr content is preferably 0.95% or less.
  • the Cr content is more preferably 0.9% or less, and even more preferably 0.85% or less.
  • Mo 0 to 0.60%
  • Mo is an element effective in improving toughness and increasing strength, and is an element effective in improving SSCC resistance and HIC resistance.
  • the Mo content may be 0% or more, but to obtain this effect, it is preferable to contain 0.01% or more. It is more preferable to contain 0.10% or more.
  • the Mo content is 0.60% or less.
  • the Mo content is preferably 0.50% or less. More preferably, it is 0.40% or less. Further preferably, it is 0.35% or less.
  • W 0 to 1.0% W contributes to increasing the strength of the steel material.
  • the W content may be 0% or more, but in order to obtain the above effect, when W is contained, the content is preferably 0.01% or more.
  • the W content exceeds 1.0%, the effect is saturated and becomes a factor of increasing costs, so when W is contained, the W content is 1.0% or less.
  • the W content is preferably 0.9% or less, and more preferably 0.8% or less. In order to suppress costs, it is even more preferable to make it 0.5% or less.
  • V is an element that can be optionally contained to increase the strength and toughness of the steel material.
  • the V content may be 0% or more, but if the content is less than 0.01%, the effect of the inclusion is poor, so the V content is preferably 0.01% or more.
  • the V content is more preferably 0.03% or more.
  • the toughness of the weld deteriorates, so if it is contained, it is preferably 0.10% or less.
  • the V content is preferably 0.09% or less.
  • the V content is more preferably 0.07% or less, and even more preferably 0.06% or less.
  • Zr, Mg and REM are elements that can be added at will to improve toughness through grain refinement and crack resistance through control of inclusion properties.
  • the content may be 0% or more, but since the effect of containing them is poor when the content is less than 0.0001%, the content is preferably 0.0001% or more. More preferably, it is 0.0005% or more.
  • the Zr content is preferably 0.0001% or more.
  • the Zr content is more preferably 0.0005% or more.
  • the REM content is preferably 0.0001% or more.
  • the REM content is more preferably 0.0005% or more.
  • the Mg content is preferably 0.0001% or more.
  • the Mg content is more preferably 0.0005% or more.
  • the Zr content should be 0.050% or less. It is preferable that the Zr content be 0.040% or less. It is more preferable that the Zr content be 0.020% or less. If they are contained, the REM content should be 0.01% or less. It is preferable that the REM content be 0.009% or less. It is more preferable that the REM content be 0.008% or less. If they are contained, the Mg content should be 0.01% or less. It is preferable that the Mg content be 0.009% or less. It is more preferable that the Mg content be 0.008% or less.
  • B 0 to 0.0020%
  • B is an element that improves hardenability, contributes to increasing the strength of the steel material, inhibits the coarsening of prior austenite grains, and improves various properties of the material.
  • the B content may be 0% or more, but in order to obtain the above effect, the content is preferably 0.0001% or more. More preferably, it is 0.0008% or more.
  • the B content exceeds 0.0020%, the effect is saturated and becomes a factor of increasing costs, so when B is contained, the B content is 0.0020% or less.
  • the B content is preferably 0.0014% or less.
  • the B content is more preferably 0.0012% or less. In order to suppress costs, it is even more preferably 0.0010% or less.
  • Ta 0 to 0.2%
  • Ta is an element that forms carbides and nitrides and contributes to improving strength.
  • the Ta content may be 0% or more, but in order to obtain the above effect, it is preferable that the Ta content is 0.0001% or more. More preferably, the Ta content is 0.0008% or more.
  • the content exceeds 0.2%, it may cause a decrease in toughness, so when Ta is contained, the Ta content is 0.2% or less. It is preferable that Ta is 0.16% or less. It is more preferable that Ta is 0.12% or less, and even more preferable that Ta is 0.10% or less.
  • Hf 0 to 0.2%
  • Re 0 to 0.005%
  • the contents of these elements are preferably 0.0001% or more.
  • they are 0.0010% or more. That is, when these elements are contained, the Hf content is preferably 0.0001% or more.
  • the Hf content is more preferably 0.0010% or more.
  • the Re content is preferably 0.0001% or more.
  • the Re content is preferably 0.001% or more.
  • Hf when these elements are contained, if the content of Hf exceeds 0.2% and the content of Re exceeds 0.005%, oxides increase and, if they aggregate, hydrogen resistance properties are impaired, so Hf is 0.2% or less and the Re content is 0.005% or less. That is, when these elements are contained, the Hf content is 0.2% or less.
  • the Hf content is preferably 0.18% or less, and more preferably 0.12% or less.
  • the Re content is 0.005% or less.
  • the Re content is preferably 0.004% or less, and more preferably 0.003% or less.
  • Sn and Sb 0 to 0.3%
  • Sn and Sb contents may be 0% or more, but in order to obtain the above effects, it is preferable that the contents are each 0.0001% or more. Preferably, they are 0.001% or more. That is, when Sn is contained, the Sn content may be 0% or more, but it is preferable that the Sn content is 0.0001% or more. It is more preferable that the Sn content is 0.001% or more. When Sb is contained, the Sb content may be 0% or more, but it is preferable that the Sb content is 0.0001% or more.
  • the Sb content is 0.001% or more.
  • the Sn content is 0.3% or less. It is preferable to make the Sn content 0.2% or less. It is more preferable to make the Sn content 0.1% or less. It is even more preferable to make the Sn content 0.01% or less.
  • Sb is contained, the Sb content is 0.3% or less. It is preferable to make the Sb content 0.2% or less. It is more preferable to make the Sb content 0.1% or less. It is even more preferable to make the Sb content 0.01% or less.
  • the Sb content is 0.3% or less. It is preferable to make the Sb content 0.2% or less. It is more preferable to make the Sb content 0.1% or less. It is even more preferable to make the Sb content 0.01% or less.
  • the remainder other than the above-mentioned components consists of Fe and unavoidable impurity elements.
  • the metal structure of the steel material of the present invention is described below.
  • Metal structure 15 inclusions/ 100 mm2 or less with an aspect ratio of 2.0 or more and a length of 10 ⁇ m or more
  • inclusions in the material include elongated MnS and cementite. These act as hydrogen accumulation sources, leading to a significant decrease in HIC resistance and causing a decrease in the hydrogen-induced crack propagation limit K IH .
  • the number of inclusions with an aspect ratio of 2.0 or more and a length of 10 ⁇ m or more is set to 15 inclusions/100 mm2 or less.
  • the number density of the inclusions is preferably 10 inclusions/100 mm2 or less.
  • the lower limit is not particularly limited, and may be 0 inclusions/100 mm2 .
  • Retained austenite is 0-3% (preferable) When retained austenite remains in the steel structure, it acts as a hydrogen trap site, increasing the amount of hydrogen in the steel and increasing the hydrogen embrittlement susceptibility. Furthermore, when a steel material or steel pipe is used as a steel structure, if the retained austenite transforms into martensite due to stress load during use, the martensite is very hard and may become a source or propagation path of HIC, significantly decreasing K IH . In the present invention, the retained austenite is set to 3% or less to improve K IH . For this reason, the retained austenite is preferably set to 3% or less. The retained austenite is more preferably set to 2% or less. More preferably, it is set to 1% or less. The retained austenite may be 0%.
  • the area fraction of bainite in the range from the steel surface (in the case of steel pipes, the surface of the steel pipe inner surface) to the center of the plate thickness is 90% or more (preferred).
  • the steel material In order to achieve a high strength of tensile strength of 520 MPa or more as a material suitable for line pipes, the steel material must have a bainite structure.
  • the bainite structure includes bainitic ferrite or granular bainite that transforms during or after accelerated cooling, which contributes to transformation strengthening, and also includes tempered bainite.
  • bainite structure includes heterogeneous structures such as ferrite, martensite, pearlite, island martensite, and retained austenite, the strength decreases, and the normal (atmospheric) toughness and K IH deteriorate. Furthermore, the presence of steel structures with different hardnesses causes stress distribution in the steel material when stress is applied during use, and acts as a hydrogen accumulation source due to stress-induced diffusion, thereby degrading HIC resistance. For this reason, it is preferable that bainite is 90% or more in area fraction. It is more preferable that bainite is 92% or more in area fraction, and even more preferable that bainite is 95% or more in area fraction. There is no particular upper limit, but it may be 100%.
  • the maximum grain size in the range from the steel surface (the surface of the steel pipe inner surface in the case of a steel pipe) to the center of the plate thickness is 25 ⁇ m or less.
  • grains with a maximum grain size of more than 25 ⁇ m from the surface of the inner steel material to the center of the plate thickness tend to accumulate strain around the grains, which easily become the starting point and propagation path of hydrogen cracks, and therefore the K IH is significantly deteriorated. Therefore, it is necessary to make the maximum grain size from the inner surface of the steel material to the center of the plate thickness 25 ⁇ m or less.
  • the maximum grain size from the inner surface of the steel material to the center of the plate thickness is preferably 24 ⁇ m or less, more preferably 22 ⁇ m or less, and even more preferably 20 ⁇ m or less. Although the lower limit is not particularly limited, the maximum grain size is preferably 4 ⁇ m or more.
  • the measurement range of the crystal grain size is 1 mm ⁇ 1 mm, and the crystal grain size is defined as the area grain size (weighted average when the boundary with an orientation difference of 15° or more is defined as the grain boundary).
  • the hydrogen-induced crack propagation threshold K IH in a high-pressure hydrogen gas environment of 1 MPa or more is 80 MPa ⁇ m 1/2 or more.
  • the high-strength steel material of the present disclosure has a hydrogen-induced crack propagation threshold K IH of 80 MPa ⁇ m 1/2 or more in a high-pressure hydrogen gas environment of 1 MPa or more.
  • the upper limit is not particularly limited, it is preferable that the hydrogen-induced crack propagation threshold K IH of the steel material is 120 MPa ⁇ m 1/2 or less, and more preferably 100 Pa ⁇ m 1/2 or less.
  • the hydrogen-induced crack propagation threshold K IH refers to the plane strain fracture toughness K IC or its provisional value obtained in a high-pressure hydrogen gas of 1 MPa or more in accordance with ASTM E399 or ASTM E1820, or the crack propagation threshold or its provisional value obtained in accordance with ASTM E1681.
  • the thickness of the steel plate is not particularly limited, but it is preferable that the thickness be 5 mm or more. It is preferable that the thickness be 30 mm or less.
  • the present invention by having the above-mentioned chemical composition and metal structure, can obtain an excellent hydrogen-induced cracking threshold K IH in high-pressure hydrogen gas, and can be applied to hydrogen line pipes.
  • the high-strength steel for line pipes according to the present invention can be obtained by limiting the manufacturing conditions shown below, and the manufacturing method and conditions are specifically explained below.
  • molten steel process Average cooling rate of molten steel: 50° C./min or more (preferred conditions)
  • the inclusions defined in the present invention aggregate during the cooling process of the molten steel, it is also effective to increase the average cooling rate of the molten steel.
  • Heating process [heating temperature of cast piece: 1000 to 1250°C] If the heating temperature of the billet or slab is less than 1000°C, the diffusion of micro-segregated impurity elements such as C, P, and S is insufficient, and a homogeneous material cannot be obtained, which causes an increase in the number of inclusions and non-uniform precipitation, thereby reducing toughness. Therefore, the heating temperature of the billet is set to 1000°C or higher.
  • the heating temperature of the billet is preferably set to 1050°C or higher, and more preferably set to 1100°C or higher.
  • the heating temperature of the billet is set to 1250°C or lower.
  • the heating temperature of the billet is preferably set to 1200°C or lower, and more preferably set to 1150°C or lower.
  • Total rolling reduction in the recrystallization temperature range after heating of the slab 35% to 55%
  • the total reduction in the recrystallization temperature range is set to 35% or more.
  • the total reduction in the recrystallization temperature range is more preferably set to 40% or more, and even more preferably set to 43% or more.
  • the total reduction in the recrystallization temperature range is set to 55% or less. Preferably, it is set to 52% or less.
  • the total reduction in the recrystallization temperature range is more preferably set to 50% or less, and even more preferably set to 48% or less.
  • the lower limit temperature Tnr of recrystallization can be calculated, for example, from the components of the steel by the following formula.
  • the surface temperature of the steel sheet can be measured by a radiation thermometer, etc.
  • [%X] indicates the content (mass%) of the X element in the steel.
  • the reduction in the final rolling pass in the recrystallization temperature range is preferably 11% or more.
  • the reduction in the final rolling pass in the recrystallization temperature range is more preferably 13% or more, and even more preferably 15% or more.
  • the reduction ratio of the final rolling pass at (recrystallization temperature - 80°C) or higher is set to 15% or more.
  • the reduction ratio of the final rolling pass at (recrystallization temperature - 80°C) or higher is preferably set to 16% or more.
  • the reduction ratio of the final rolling pass at (recrystallization temperature - 80°C) or higher is more preferably 18% or more, and even more preferably 20% or more.
  • Rolling at a temperature lower than (recrystallization temperature - 80°C) is more effective at refining the grains because more strain is introduced at lower temperatures. For this reason, it is preferable to roll at low temperatures within the range where the cooling start temperature for controlled cooling can be observed.
  • the lower the rolling end temperature the better.
  • the Ar 3 point means the ferrite transformation start temperature during cooling, and can be calculated, for example, from the composition of the steel by the following formula.
  • the surface temperature of the hot-rolled steel sheet can be measured by a radiation thermometer or the like.
  • Ar3 (°C) 910-310[%C]-80[%Mn]-20[%Cu]-15[%Cr]-55[%Ni]-80[%Mo]
  • [%X] indicates the content (mass%) of the X element in the steel.
  • Controlled cooling start temperature Surface temperature of hot-rolled steel sheet above Ar3 transformation point
  • Ar3 point Ar3 transformation point
  • ferrite is generated before controlled cooling, and the strength is greatly reduced. Therefore, the surface temperature of the hot-rolled steel sheet at the start of cooling is set to be equal to or higher than the Ar3 transformation point.
  • the surface temperature of the hot-rolled steel sheet at the start of cooling is preferably equal to or higher than the Ar3 transformation point + 20°C, and more preferably equal to or higher than the Ar3 transformation point + 50°C.
  • the surface temperature of the hot-rolled steel sheet at the start of cooling is the temperature of the tail end of the hot-rolled steel sheet, where the cooling start temperature is the lowest.
  • the surface temperature of the hot-rolled steel sheet at the start of cooling is preferably equal to or lower than the Ar3 transformation point + 120°C, and more preferably equal to or lower than the Ar3 transformation point + 80°C.
  • the cooling start time difference between the leading end and the tail end of the hot-rolled steel sheet is set to 50 seconds or less.
  • the cooling start time difference is preferably set to 45 seconds or less.
  • the cooling start time difference is more preferably set to 40 seconds or less, and further preferably set to 32 seconds or less. Although it is possible to shorten the cooling start time difference by shortening the length of the hot-rolled steel sheet, this reduces manufacturability, so it is preferable to shorten the cooling start time difference by increasing the hot-rolled steel sheet conveying speed.
  • the cooling start time difference may be 0 seconds, but from the viewpoint of manufacturability, it is preferable to set it to 20 seconds or more.
  • Average cooling rate from 750 ° C to 550 ° C at the center of plate thickness 15 to 50 ° C / s If the average cooling rate from 750°C to 550°C at the center of the plate thickness is less than 15°C/s, the specified bainite structure including granular bainite cannot be obtained, and strength is reduced. For this reason, the average cooling rate at the center of the plate thickness is set to 15°C/s or more. From the viewpoint of suppressing the variation in the structure, the average cooling rate at the center of the plate thickness is preferably set to 17°C/s or more.
  • the average cooling rate at the center of the plate thickness is preferably set to 20°C/s or more, and more preferably set to 25°C/s or more.
  • the average cooling rate is set to 50°C/s or less.
  • the average cooling rate is preferably set to 48°C/s or less, and more preferably set to 45°C/s or less.
  • the average cooling rate is more preferably set to 42°C/s or less, and most preferably set to 38°C/s or less.
  • the cooling to 550°C or less at the hot-rolled steel plate temperature at the center of the plate thickness is not particularly limited, but from the viewpoint of suppressing the variation in the structure and grain size, the average cooling rate is preferably set to 15°C/s or more and 50°C/s or less. That is, for cooling to 550°C or less, the average cooling rate is preferably 15°C/s or more. The average cooling rate is more preferably 30°C/s or more, and even more preferably 35°C/s or more. For cooling to 550°C or less, the average cooling rate is preferably 50°C/s or less. The average cooling rate is more preferably 48°C/s or less, and even more preferably 42°C/s or less. The average cooling rate for 550°C or less is the average value of the cooling rates from 550°C to 250°C.
  • the cooling stop temperature at the center of the thickness after hot rolling exceeds 650°C, the material strength is significantly reduced, and from the viewpoint of obtaining a uniform bainite structure, the cooling stop temperature at the center of the thickness is set to 650°C or less.
  • the cooling stop temperature at the center of the thickness is preferably set to 620°C or less, more preferably set to 615°C or less, and even more preferably set to 600°C or less.
  • the cooling stop temperature at the center of the thickness is less than 250°C, quench cracks are likely to occur during cooling.
  • the cooling stop temperature is set to 250°C or more.
  • the cooling stop temperature at the center of the thickness is preferably set to 300°C or more, more preferably set to 350°C or more, and even more preferably set to 380°C or more.
  • the cooling stop temperature needs to be set to a predetermined temperature or more. Specifically, hydrogen present in the steel gradually escapes during cooling, and the higher the temperature, the greater the effect, but if the cooling stop temperature is too low, the steel will be overcooled and hydrogen will remain in the steel. Furthermore, if the cooling stop temperature is too low, it is easy to form retained austenite, which rapidly increases hydrogen in a large amount compared to other phases. Therefore, in order to reduce the amount of hydrogen in the steel, the cooling stop temperature needs to be 250°C or higher. Although it is acceptable to allow the steel to cool after cooling is stopped, it is more preferable to cool it slowly until the temperature drops by about 50°C from the cooling stop temperature in order to promote the formation of bainite.
  • Dehydrogenation treatment (optimal conditions) If hydrogen is present in steel to begin with, the acceleration of fatigue crack growth increases, and the fatigue life decreases. Therefore, it is preferable to use a dehydrogenation process to release the hydrogen remaining after manufacturing. Dehydrogenation can reduce the amount of hydrogen in steel by holding the product at high temperature for a certain period of time before use. Hydrogen can also be dehydrogenated by holding it at room temperature for a long time. When holding it at room temperature, the holding time is long, so the holding time is preferably 96 hours or more. Furthermore, since scale on the steel surface inhibits dehydrogenation, it is preferable to remove the scale before dehydrogenation.
  • the holding time R (sec) is preferably calculated from the plate thickness and pipe thickness t (mm) of the steel plate and steel pipe, and the hydrogen diffusion coefficient D (mm ⁇ sec ⁇ 1 ) in steel at room temperature, as shown in the following formula (A).
  • R ⁇ t2 /D (A) The hydrogen diffusion coefficient varies depending on the contained components and metal structure, but for example, the hydrogen diffusion coefficient may be 1 ⁇ 10 ⁇ 5 to 5 ⁇ 10 ⁇ 3 mm 2 /s, and more preferably 5 ⁇ 10 ⁇ 4 mm 2 /s or less.
  • the dehydrogenation process is carried out before pipe making or welding to connect steel pipes. It is preferable to carry out the dehydrogenation process at a high temperature because the hydrogen diffusion coefficient D at high temperatures becomes small and hydrogen is quickly removed.
  • the diffusion coefficient D' (diffusion coefficient at each temperature) at which the value of D in the above formula (A) is maintained may be used for calculation.
  • the dehydrogenation temperature is preferably 550°C or less. It is more preferable that the dehydrogenation temperature T is 500°C or less. It is even more preferable that the dehydrogenation temperature T is 400°C or less, and most preferably 300°C or less.
  • the dehydrogenation temperature T is room temperature or higher.
  • the dehydrogenation temperature T is 50°C or higher. It is more preferable that the dehydrogenation temperature T is 100°C or higher, and most preferably 150°C or higher.
  • the dehydrogenation temperature T mentioned here is the temperature of the atmosphere in the dehydrogenation process. Room temperature means 20 ⁇ 10°C.
  • the temperature Tc at the dehydrogenation temperature T (atmospheric temperature) for R (sec) or more as specified by formula (A), and it is even more preferable to hold the temperature Tc for the above-mentioned holding time R (sec) or more after it reaches the target dehydrogenation temperature T.
  • at least the former can appropriately control the amount of hydrogen in the surface layer of the steel material and steel pipe, and by implementing the latter, the amount of hydrogen in the steel material from the surface layer to the center of the thickness of the steel material and steel pipe can be appropriately controlled.
  • the thickness temperature, or the center temperature Tc can be measured using a thermocouple or the like, or can be predicted using the finite element method or the like.
  • the time and temperature of the dehydrogenation process may include the temperature and time applied when heating in the pipe-making process for electric resistance welded pipes, UOE, etc., as described below.
  • scale on the steel surface inhibits dehydrogenation, it is preferable to remove the scale before carrying out the dehydrogenation process.
  • the removal method There is no restriction on the removal method, but it may be physical cleaning using a high-pressure cleaner, for example, or a chemical method using a scale remover. The effect of scale removal can be obtained if a thickness of about 100 ⁇ m is removed.
  • a UOE steel pipe which is an example of a steel pipe for high strength line pipe, can be obtained by limiting the manufacturing conditions shown below, and the manufacturing method and conditions will be specifically described.
  • the chemical composition, metal structure, and hydrogen-induced crack propagation lower limit K IH of the UOE steel pipe are the same as those described for the steel plate of the first embodiment, and as for the manufacturing method, the molten steel process, heating process, hot rolling process, controlled cooling process after hot rolling, and dehydrogenation process are performed in the same manner as those described for the steel material.
  • the pipe-making process after rolling will be specifically described below.
  • UOE steel pipes are manufactured by bending hot-rolled steel sheets, specifically by groove-forming the ends of the hot-rolled steel sheets, and forming them into a steel pipe shape using a C press, a U press, and an O press, then seam-welding the butt joints using internal and external welding, and then expanding the pipe as necessary.
  • Any welding method may be used as long as it provides sufficient joint strength and joint toughness, but it is preferable to use submerged arc welding from the viewpoint of excellent welding quality and manufacturing efficiency.
  • Pipe expansion can also be performed on steel pipes that have been formed into a tubular shape by press bending and then have seam-welded butt joints.
  • the average cooling rate of the welded steel pipe in the temperature range from 1500°C to 1000°C to 50°C/min or more.
  • the average cooling rate is more preferably 55° C./min or more, and even more preferably 60° C./min or more. Although there is no particular upper limit, the average cooling rate is preferably 100° C./min or less.
  • an example of a steel pipe for high strength line pipe according to the present invention is an electric resistance welded steel pipe, which can be obtained by limiting the manufacturing conditions shown below, and the manufacturing method and conditions will be specifically described below.
  • the composition, metal structure, and hydrogen induced crack propagation lower limit K IH of the steel material are the same as those described for the steel material of the first embodiment, and as for the manufacturing method, the steps other than the cooling step after rolling and the pipe making step (melting step, heating step, hot rolling step, dehydrogenation treatment step) are performed in the same manner as those described for the steel material.
  • Cooling process after rolling (controlled cooling process)
  • the cooling start temperature and the average cooling rate of the controlled cooling are the same as those described in the first embodiment.
  • the cooling stop temperature at the center of the thickness after hot rolling exceeds 650°C, the material strength is significantly reduced, and from the viewpoint of obtaining a uniform bainite structure, the cooling stop temperature at the center of the thickness is set to 650°C or less.
  • the cooling stop temperature at the center of the thickness is preferably set to 620°C or less, more preferably set to 615°C or less, and even more preferably set to 600°C or less.
  • the cooling stop temperature at the center of the thickness is set to 250°C or more.
  • the cooling stop temperature at the center of the thickness is preferably set to 300°C or more, more preferably set to 350°C or more, and even more preferably set to 380°C or more.
  • the cooling stop temperature at the center of the thickness is most preferably 450°C or more. After cooling is stopped, it is sufficient to allow the steel to cool, but in order to promote the generation of bainite, it is more preferable to slowly cool the steel until the temperature drops by about 50°C from the cooling stop temperature.
  • the winding temperature is preferably 650°C or less.
  • the winding temperature is more preferably 620°C or less, more preferably 615°C or less, and even more preferably 600°C or less.
  • the winding temperature is preferably 250°C or more, more preferably 300°C or more, more preferably 350°C or more, and most preferably 380°C or more.
  • the electric resistance welded steel pipe given as an example of the present invention is manufactured by forming the pipe into a cylindrical shape by cold roll forming, and then butting and welding both circumferential ends of the cylindrical shape together. Furthermore, the electric resistance welded steel pipe may be manufactured by forming the pipe into an electric resistance welded steel pipe material (electric resistance welded steel pipe) using a sizing roll that satisfies the following formula (1) (sizing process), and applying an internal pressure p (MPa) that satisfies the following formula (2) to the inner surface of the electric resistance welded steel pipe material (internal pressure application process).
  • the cylindrical shape means that the circumferential cross section of the tube is in a "C" shape.
  • the plate thickness of the hot-rolled steel plate means the plate thickness of the hot-rolled steel plate before the sizing process is performed.
  • X thickness of electric welded steel pipe material (mm) / radius of electric welded steel pipe material (mm)) ⁇ yield strength of electric welded steel pipe material (MPa)
  • MPa yield strength of electric welded steel pipe material
  • the thickness of the electric welded steel pipe material given as an example of the steel pipe of the present invention is preferably 5 mm or more.
  • the thickness of the electric welded steel pipe material is preferably 30 mm or less.
  • the radius of the electric welded pipe material is preferably 400 mm or less.
  • the radius of the electric welded pipe material is preferably 200 mm or more.
  • the yield strength of the electric welded steel pipe material is preferably 480 MPa or more in order to withstand the gas pressure of pipeline operation. A yield strength of 500 MPa or more is more preferable.
  • the yield strength is preferably 560 MPa or less.
  • a yield strength of 550 MPa or less is more preferable.
  • 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 axial direction of the tube. If the diameter of the sizing roll is less than the right side of the formula (1), the intended residual shear stress of the present invention cannot be obtained.
  • the diameter of the sizing roll is preferably 2000 mm or less.
  • the electric resistance welded steel pipe material is expanded to generate tensile stress in the circumferential direction of the pipe, thereby reducing the absolute value of the residual stress in the circumferential direction of the pipe.
  • the left side (X) of the above equation (2) corresponds to the internal pressure p when the tensile stress generated in the circumferential direction of the pipe is equal to the yield stress of the electric resistance welded steel pipe material.
  • the internal pressure p is set to a value greater than the left side (X) of equation (2) and the electric resistance welded steel pipe material is expanded to the plastic region.
  • the high-strength steel material disclosed herein can be formed into a tube by press bending, roll forming, UOE forming, or the like, and then the butt joints can be welded to produce high-strength steel pipes for sour-resistant line pipes (UOE steel pipes, electric resistance welded steel pipes, spiral steel pipes, etc.) with excellent material uniformity within the steel plate, suitable for transporting crude oil or natural gas.
  • sour-resistant line pipes UOE steel pipes, electric resistance welded steel pipes, spiral steel pipes, etc.
  • steel pipes with excellent HISC resistance can be produced, even if there is a high hardness region in the weld.
  • the above-mentioned observation surface was etched using a 3 vol% nital solution, and a scanning electron microscope photograph was taken at an appropriate magnification between 1000 and 5000 times to observe bainite.
  • Bainite was judged visually by comparison with the structure photograph in Non-Patent Document 1, and the structure fraction was determined by binarizing the bainite and other regions in the SEM photograph based on the above judgment, and determining the fraction by image analysis, which was taken as the area fraction of bainite.
  • Samples for metal structure observation were taken from the center of the plate width in the longitudinal center of the steel material and steel pipe obtained as described above, and the cross section of this sample parallel to the rolling direction was used as the observation surface.
  • the observation surface was mirror-polished and then etched with colloidal silica, and observation was performed with a scanning electron microscope (SEM) at the center of the sample in a field of view of 10 mm x 10 mm. The observation was performed at a magnification of 2000 to 5000 times, and the average of three fields of view was used as the number density of inclusions.
  • SEM scanning electron microscope
  • Temperature-programmed hydrogen analysis The amount of hydrogen remaining in the steel was measured using a temperature-programmed desorption analysis method, using a low-temperature temperature-programmed hydrogen analyzer (gas chromatograph type) (JTF-20AL). Temperature-programmed desorption analysis was performed in the temperature range from room temperature to 400°C at a heating rate of 200°C/h, and the sum of the measurements was taken as the amount of hydrogen.
  • the test specimens were cylindrical, 30 mm long in the longitudinal direction of the steel pipe, at a 1/4 position of the plate thickness of the steel material and a 1/4 position from the inner surface of the steel pipe, and had a diameter of 7 ⁇ . This amount of hydrogen was measured before the steel was subjected to the high-pressure hydrogen fatigue test described in the aging section below, and is the amount of H shown in Tables 1-1 and 1-2.
  • Fracture toughness test in high pressure hydrogen gas The test was carried out in accordance with ASTM E1820 at room temperature (20 ⁇ 10°C), in hydrogen gas (containing 100% hydrogen) at a pressure of 25 MPa, or in a mixed atmosphere of natural gas (mainly hydrocarbons such as methane and ethane) containing hydrogen with a partial pressure of 1 MPa or more at the above temperature and pressure.
  • CT test pieces (plate thickness 12.7 mm, plate width 25.4 mm) were used as test pieces, and were taken in a direction in which the machine notch introduction direction and the rolling direction of the steel material were parallel. Fatigue pre-cracks were introduced in the atmosphere under the conditions of frequency: 1 Hz, repeated load waveform: sine wave, control method: K value control, and stress ratio R: 0.1.
  • the test was carried out in hydrogen gas or a mixed atmosphere of hydrogen gas and natural gas.
  • the fracture toughness test was carried out by the unloading-elastic compliance method using a single test piece.
  • the crosshead displacement speed during loading was 0.002 mm/sec.
  • Example 1 The dehydrogenation treatment of steel pipes No. 2, 4, 8, 14, 22, and 33 carried out in Example 1 was carried out at the dehydrogenation treatment temperature T (ambient temperature) and time shown in Tables 2-1 and 2-2, which correspond to the dehydrogenation holding time t in Table 4 being Y and the holding time tc at the plate thickness center temperature Tc being N, respectively.
  • the dehydrogenation temperature T was set to the temperature shown in Table 4, and the holding time tc after the center temperature Tc reached the dehydrogenation temperature T shown in Table 4 was set to satisfy formula (A).
  • the dehydrogenation temperature T is the temperature shown in Table 4, but the holding time t of the ambient temperature and the holding time tc after the temperature at the center of the plate thickness Tc reaches the above-mentioned dehydrogenation temperature T do not satisfy the above-mentioned formula (A).
  • dehydrogenation holding time t is Y
  • dehydrogenation temperature T ambient temperature
  • holding time t satisfies formula (A)
  • dehydrogenation holding time t is N
  • dehydrogenation temperature T ambient temperature
  • holding time tc at steel center temperature Tc is Y means that the holding time tc after the plate thickness center temperature Tc reaches a predetermined temperature satisfies formula (A), and "holding time tc at steel center temperature Tc is N” means that the plate thickness center temperature Tc reaches a predetermined temperature, but the holding time tc after Tc reaches the predetermined temperature does not satisfy formula (A).
  • All of the inventive examples of the present invention satisfied the conditions of a hydrogen-induced crack propagation threshold K IH of 80 MPa ⁇ m 1/2 or more and a tensile strength of 520 MPa or more.
  • the examples in which the dehydrogenation treatment was performed under more suitable conditions had superior fracture toughness in hydrogen.

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Abstract

The purpose of the present invention is to provide a high-strength line pipe steel material that has excellent fracture toughness in hydrogen in a high-pressure hydrogen gas environment, and that is suitable for use for a steel structure used in a high-pressure hydrogen gas environment, such as a line pipe for 100% hydrogen gas or for natural gas (natural gas is gas containing a hydrocarbon such as methane and ethane as the main component) containing hydrogen at a hydrogen partial pressure of 1 MPa or more; a method for manufacturing the steel material; a steel tube for high-strength line pipes; and a method for manufacturing the steel tube.　This high-strength line pipe steel material having excellent fracture toughness in hydrogen is characterized by having a specific compositional makeup and a specific structure, having a tensile strength of 520 MPa or more, and having a hydrogen induced crack propagation lower limit KIH of 80 MPa·m1/2 or more in a high-pressure hydrogen gas environment of 1 MPa or more.

Description

水素中破壊靭性に優れた高強度ラインパイプ用鋼材、その製造方法、高強度ラインパイプ用鋼管およびその製造方法High-strength line pipe steel material with excellent fracture toughness in hydrogen, manufacturing method thereof, high-strength line pipe steel pipe and manufacturing method thereof
　本発明は、水素ガスの輸送用ラインパイプ等の用途に好適な、１ＭＰａ以上の高圧水素ガス環境において水素中破壊靭性に優れた高強度ラインパイプ用鋼材、その製造方法、高強度ラインパイプ用鋼管とその製造方法に関する。 The present invention relates to a high-strength line pipe steel material that has excellent fracture toughness in hydrogen in a high-pressure hydrogen gas environment of 1 MPa or more, and is suitable for applications such as line pipes for transporting hydrogen gas, a manufacturing method thereof, and a high-strength line pipe steel pipe and a manufacturing method thereof.
　既存のエネルギーインフラとして、原油や天然ガス等を輸送するラインパイプがある。これらの鋼構造物は、硫化水素を含有する雰囲気で使用され、水素誘起割れ（ＨＩＣ）や硫化物応力腐食割れ（ＳＳＣＣ）等の水素脆化の発生が安全上の問題となっており、その抑制が求められてきた。水素誘起割れや、硫化物応力腐食割れ等の水素脆化の発生を防止するために、割れ起点となる鋼材中のＭｎＳの量の低減、およびＴｉ、Ｎｂの炭窒化物や酸化物の集積抑制、あるいは中心偏析の硬化相の偏析の抑制など、様々な対策が講じられてきた。また、鋼材の耐食性の向上による割れ起点発生の抑制という観点では、鋼材へのＳｎやＳｂの添加が提案されている（例えば、特許文献１、２）。 Existing energy infrastructure includes line pipes for transporting crude oil, natural gas, etc. These steel structures are used in an atmosphere containing hydrogen sulfide, and hydrogen embrittlement such as hydrogen induced cracking (HIC) and sulfide stress corrosion cracking (SSCC) has become a safety issue, and there has been a demand for its prevention. In order to prevent hydrogen embrittlement such as hydrogen induced cracking and sulfide stress corrosion cracking, various measures have been taken, such as reducing the amount of MnS in the steel material, which is the crack initiation point, suppressing the accumulation of Ti and Nb carbonitrides and oxides, and suppressing the segregation of the hardened phase of the central segregation. In addition, from the viewpoint of suppressing the occurrence of crack initiation points by improving the corrosion resistance of steel material, the addition of Sn and Sb to steel material has been proposed (for example, Patent Documents 1 and 2).
　近年、脱炭素社会構築を目的としたクリーンなエネルギー源として水素の活用が推進されている。そのため、水素ガスの大量輸送を目的とし、天然ガスラインパイプに水素を一定割合で混合した天然ガスや、水素ガスを代替として圧送する水素ガス輸送網の構築が検討されている。これらのパイプライン運転時の輸送圧力は、１～４０ＭＰａの高圧力が想定されており、ラインパイプは、高圧力の水素ガス環境に暴露されることになる。このような環境で使用される鋼材には、既存のサワー環境で要求される、鋼管内表面における腐食発生の抑制や材料内への水素集積を低減する特性に加え、水素ガス環境で要求される、耐水素性を兼ね備える必要がある。 In recent years, the use of hydrogen has been promoted as a clean energy source with the aim of building a decarbonized society. As a result, the construction of a hydrogen gas transportation network that uses natural gas mixed with a certain ratio of hydrogen in natural gas line pipes and pressurized hydrogen gas as an alternative for the purpose of transporting large amounts of hydrogen gas is being considered. The transport pressure during operation of these pipelines is expected to be high, at 1 to 40 MPa, and the line pipes will be exposed to a high-pressure hydrogen gas environment. Steel materials used in such environments need to have the hydrogen resistance required in hydrogen gas environments, in addition to the properties required in existing sour environments, such as inhibiting the occurrence of corrosion on the inner surface of steel pipes and reducing hydrogen accumulation within the material.
　高圧水素ガス環境下で使用される鋼構造物には、水素中破壊靭性を示す、ＳＵＳ３１６Ｌ等のオーステナイト系ステンレス鋼が使用されている。しかし、鋼材のコストが高いことに加え、低強度であり、高い水素圧に耐えうるように設計した場合、肉厚が厚くなるとともにラインパイプの価格は非常に高価となり、パイプライン敷設には適さない。そのため、水素用ラインパイプ向けとして、より低コストで、かつ高圧水素ガス環境にも耐えうる鋼材が要望されてきた。 Austenitic stainless steels such as SUS316L, which exhibit fracture toughness in hydrogen, are used for steel structures used in high-pressure hydrogen gas environments. However, in addition to being expensive, the steel has low strength, and when designed to withstand high hydrogen pressure, the wall thickness becomes thicker and the line pipe becomes very expensive, making it unsuitable for pipeline construction. For this reason, there has been a demand for steel materials for hydrogen line pipes that are lower cost and can withstand high-pressure hydrogen gas environments.
　上記の問題を解決するために、例えば特許文献３には、Ｍｎの含有量が多いオーステナイト系鋼材が提案されている。特許文献３に記載の技術によって、オーステナイト系ステンレス鋼と比較し、低コストである鋼材の提供が可能であるが、オーステナイト系であるため、低合金鋼と比べると高コストである。また、耐ＨＩＣ性や、耐ＳＳＣＣ性など、水素誘起き裂の起点となる孔食の抑制は考慮されていない。 In order to solve the above problems, for example, Patent Document 3 proposes an austenitic steel material with a high Mn content. The technology described in Patent Document 3 makes it possible to provide a steel material that is less expensive than austenitic stainless steel, but because it is austenitic, it is more expensive than low-alloy steel. In addition, no consideration is given to suppressing pitting corrosion, which is the starting point of hydrogen-induced cracking, such as HIC resistance and SSCC resistance.
　また、パイプラインでは、操業開始、およびシャットダウンを繰り返し行うため、ラインパイプには、繰返し応力が負荷される。そのため、パイプラインのような鋼構造物を設計する際には、疲労破壊を考慮することが必須となる。高圧水素ガス環境中で使用する鋼構造物の疲労破壊の限界点は、パイプラインの操業条件と、鋼材の水素ガス中の破壊靭性値に相当する、水素誘起き裂進展下限界ＫＩＨから計算した限界き裂長さに対応する。水素用構造物の長寿命化、安全性向上という観点から、鋼材のＫＩＨを高くすることが一つの有効な指針とされている。 In addition, in a pipeline, the start-up and shutdown of operation are repeated, so that the line pipe is subjected to repeated stress. Therefore, when designing a steel structure such as a pipeline, it is essential to consider fatigue failure. The limit of fatigue failure of a steel structure used in a high-pressure hydrogen gas environment corresponds to the critical crack length calculated from the operating conditions of the pipeline and the hydrogen-induced crack propagation lower limit K IH , which corresponds to the fracture toughness value of the steel in hydrogen gas. From the viewpoint of extending the life and improving the safety of hydrogen structures, increasing the K IH of the steel is considered to be one effective guideline.
　水素パイプラインは、溶接部、すなわち溶接金属部、熱影響部を有するラインパイプの使用が想定されている。特許文献４では、ＫＩＨに優れた鋼材の製造方法を提案しているが、溶接部の特性には言及していない。一般的に、母材と比較し、溶接部は水素による特性劣化の影響を受けやすい。したがって、溶接部も含めたＫＩＨの向上が重要である。 Hydrogen pipelines are expected to use line pipes having welds, i.e., weld metal parts and heat-affected parts. Patent Document 4 proposes a manufacturing method for steel materials with excellent K IH , but does not mention the characteristics of the welds. In general, welds are more susceptible to property deterioration due to hydrogen than base materials. Therefore, it is important to improve K IH , including welds.
　鋼材のＫＩＨを高くするためには、例えば、粗大な炭化物を含む上部ベイナイトを低減させた方がよい。 In order to increase the K IH of a steel material, for example, it is better to reduce upper bainite containing coarse carbides.
特開２０１１－２６６９５号公報JP 2011-26695 A 特開２０１０－２０９４６１号公報JP 2010-209461 A 特表２０１９－５０５６７５号公報JP 2019-505675 A 国際公開第２０１７／０４７０９９号International Publication No. 2017/047099
　本発明は、上記した従来の問題を解決せんとしてなされたもので、１００％水素ガスまたは水素分圧が１ＭＰａ以上の水素を含む天然ガス（天然ガスはメタン、エタンなどの炭化水素を主な成分とするガス）用ラインパイプ等の、高圧水素ガス環境下で使用される鋼構造物用として好適な、高圧水素ガス環境下における水素中破壊靭性に優れた高強度ラインパイプ用鋼材、その製造方法、高強度ラインパイプ用鋼管およびその製造方法を提供することを目的とする。高圧水素ガス環境としては、１ＭＰａ以上の高圧水素ガス、もしくは水素ガスを０．２％以上含む環境が想定される。 The present invention has been made to solve the above-mentioned problems of the prior art, and aims to provide a high-strength line pipe steel material with excellent fracture toughness in hydrogen in a high-pressure hydrogen gas environment, suitable for steel structures used in a high-pressure hydrogen gas environment, such as line pipes for 100% hydrogen gas or natural gas (natural gas is a gas whose main components are hydrocarbons such as methane and ethane) containing hydrogen at a hydrogen partial pressure of 1 MPa or more, a manufacturing method thereof, and a high-strength line pipe steel pipe and a manufacturing method thereof. The high-pressure hydrogen gas environment is assumed to be high-pressure hydrogen gas of 1 MPa or more, or an environment containing 0.2% or more hydrogen gas.
　なお、ここでいう「高圧水素ガス環境下における水素中破壊靭性に優れた」とは、室温（２０±１０℃）、圧力１ＭＰａ以上の水素ガス、または水素分圧として１ＭＰａ以上の水素を含む天然ガス（主成分はメタン、エタンなどの炭化水素）混合雰囲気の両環境下で破壊靭性試験を実施して求めた、水素誘起き裂進展下限界ＫＩＨが８０ＭＰａ・ｍ１／２以上である場合をいうものとする。なお、破壊靭性値は、ＡＳＴＭ　Ｅ３９９、ＡＳＴＭ　Ｅ１８２０、およびＡＳＴＭ　Ｅ１６８１に準拠した破壊靭性試験を実施して求めたものをいう。水素分圧として１ＭＰａ以上の水素を含む天然ガスとは、例えば水素濃度が体積分率で３０％以下であり、ガス全体の圧力が３０ＭＰａ以下であるものをさす。 In addition, the term "excellent fracture toughness in hydrogen under high pressure hydrogen gas environment" refers to a case where the hydrogen-induced crack propagation lower limit K IH is 80 MPa·m 1/2 or more, which is determined by conducting a fracture toughness test under both environments of room temperature (20± 10 °C), hydrogen gas at a pressure of 1 MPa or more, or a mixed atmosphere of natural gas (mainly composed of hydrocarbons such as methane and ethane) containing hydrogen at a hydrogen partial pressure of 1 MPa or more. The fracture toughness value refers to a value determined by conducting a fracture toughness test in accordance with ASTM E399, ASTM E1820, and ASTM E1681. 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.
　ここでいう「鋼材」には、薄鋼板、厚鋼板、継目無鋼管、電縫溶接鋼管、形鋼、棒鋼等が含まれる。 The term "steel" here includes thin steel plates, thick steel plates, seamless steel pipes, electric resistance welded steel pipes, steel sections, steel bars, etc.
　本発明者らは、水素脆化の根本要因である、鋼材への水素吸収を抑制することを目的とし、高圧水素ガス環境下において水素中破壊靭性に優れた高強度ラインパイプ用鋼材及び高強度ラインパイプ用鋼管を得るための鋼材が満足すべき条件について技術検討を行った。その結果、アスペクト比が２．０以上かつ長さが１０μｍ以上の介在物が１５個／１００ｍｍ２以下であり、鋼材および鋼管表面から板厚中央の範囲における上記ベイナイトの最大粒径が２５μｍ以下である金属組織では、鋼材および鋼管の水素誘起き裂進展下限界ＫＩＨが向上することを知見した。加えて、残留オーステナイトの面積分率が０～３％であり、鋼材および鋼管表面から板厚中央の範囲における、ベイナイトが面積分率で９０％以上であれば、鋼材の水素誘起き裂進展下限界ＫＩＨがさらに向上することを知見した。このような鋼組織を実現するためには、熱間圧延工程での圧延条件および圧延後の冷却条件を厳密にコントロールする必要があり、その条件を見出すことに成功した。本発明は、これらの知見に基づいてなされたものである。また、本発明において高強度とは５２０ＭＰａ以上の引張強さを指すものとする。 The present inventors conducted technical studies on the conditions that should be satisfied by a steel material for obtaining a high-strength linepipe steel material and a high-strength linepipe steel pipe having excellent fracture toughness in hydrogen under a high-pressure hydrogen gas environment, with the aim of suppressing hydrogen absorption into the steel material, which is the root cause of hydrogen embrittlement. As a result, it was found that the hydrogen-induced crack propagation lower limit K IH of the steel material and the steel pipe is improved in a metal structure in which the number of inclusions having an aspect ratio of 2.0 or more and a length of 10 μm or more is 15 pieces/100 mm2 or less, and the maximum grain size of the bainite in the range from the surface of the steel material and the steel pipe to the center of the plate thickness is 25 μm or less. In addition, it was found that the hydrogen-induced crack propagation lower limit K IH of the steel material is further improved if the area fraction of the retained austenite is 0 to 3%, and the area fraction of the bainite in the range from the surface of the steel material and the steel pipe to the center of the plate thickness is 90% or more. In order to realize such a steel structure, it is necessary to strictly control the rolling conditions in the hot rolling process and the cooling conditions after rolling, and the inventors succeeded in finding the conditions. The present invention has been made based on these findings. In the present invention, high strength refers to a tensile strength of 520 MPa or more.
　すなわち、本発明の要旨は、以下の通りである。
［１］　質量％で、
Ｃ：０．０２～０．１５％、
Ｓｉ：０．０１～２．０％、
Ｍｎ：０．５～１．５％、
Ｐ：０．０００１～０．０１５％、
Ｓ：０．０００２～０．００１５％、
Ａｌ：０．００５～０．１５％、
Ｏ：０．０１％以下、
Ｎ：０．０１０％以下、
Ｎｂ：０．１０％以下、
Ｈ：０．０２ｐｐｍ以下を含み、
あるいはさらに、
Ｃａ：０～０．００５％、
Ｎｉ：０～２．０％、
Ｔｉ：０～０．１％、
Ｃｕ：０～１．０％、
Ｃｒ：０～１．０％、
Ｍｏ：０～０．６０％、
Ｗ：０～１．０％、
Ｖ：０～０．１０％、
Ｚｒ：０～０．０５０％、
Ｍｇ：０～０．０１％、
ＲＥＭ：０～０．０１％、
Ｂ：０～０．００２０％、
Ｔａ：０～０．２％、
Ｈｆ：０～０．２％、
Ｒｅ：０～０．００５％、
Ｓｎ：０～０．３％、
Ｓｂ：０～０．３％から選択される１種以上を含み、
残部がＦｅおよび不可避的不純物元素である、化学組成を有し、
ベイナイトおよびアスペクト比が２．０以上かつ長さが１０μｍ以上の介在物が１５個／１００ｍｍ２以下である金属組織を有し、
鋼材表面から板厚中央の範囲における前記ベイナイトの最大粒径が２５μｍ以下であり、
引張強度が５２０ＭＰａ以上であって、
１ＭＰａ以上の高圧水素ガス環境において、水素誘起き裂進展下限界ＫＩＨが８０ＭＰａ・ｍ１／２以上である水素中破壊靭性に優れた高強度ラインパイプ用鋼材。
［２］　さらに、前記化学組成が、質量％で、
Ｃａ：０．０００１～０．００５％、
Ｎｉ：０．０１～２．０％、
Ｔｉ：０．００５～０．１％、
Ｃｕ：０．０１～１．０％、
Ｃｒ：０．０１～１．０％、
Ｍｏ：０．０１～０．６０％、
Ｗ：０．０１～１．０％、
Ｖ：０．０１～０．１０％、
Ｚｒ：０．０００１～０．０５０％、
Ｍｇ：０．０００１～０．０１％、
ＲＥＭ：０．０００１～０．０１％、
Ｂ：０．０００１～０．００２０％、
Ｔａ：０．０００１～０．２％、
Ｈｆ：０．０００１～０．２％、
Ｒｅ：０．０００１～０．００５％、
Ｓｎ：０．０００１～０．３％、
Ｓｂ：０．０００１～０．３％である［１］に記載の水素中破壊靭性に優れた高強度ラインパイプ用鋼材。
［３］　残留オーステナイトが面積分率で０～３％であり、鋼材表面から板厚中央の範囲における前記ベイナイトが面積分率で９０％以上である［１］または［２］に記載の水素中破壊靭性に優れた高強度ラインパイプ用鋼材。
［４］　前記［１］または［２］に記載の成分組成を有する鋳片を１０００～１２５０℃で加熱する加熱工程と、
前記加熱工程で加熱された前記鋳片を、再結晶温度域での総圧下率が３５％以上５５％以下、かつ前記再結晶温度域での最終圧延パスの圧下率が１０％以上、かつ（再結晶温度－８０℃）以上における最終圧延パスの圧下率が１５％以上、さらに鋼板表面温度で圧延終了温度がＡｒ３変態点以上の条件で圧延する熱間圧延工程と、
前記熱間圧延工程で得られた熱延鋼板を、冷却開始温度が前記熱延鋼板の表面温度でＡｒ３変態点以上、前記熱延鋼板の先端と尾端の冷却開始時間差が５０秒以内、７５０℃から５５０℃までの平均冷却速度が板厚中央温度で１５～５０℃／ｓ、冷却停止温度が２５０～６５０℃である条件で冷却する制御冷却工程と、
を有する水素中破壊靭性に優れた高強度ラインパイプ用鋼材の製造方法。
［５］　質量％で、
Ｃ：０．０２～０．１５％、
Ｓｉ：０．０１～２．０％、
Ｍｎ：０．５～１．５％、
Ｐ：０．０００１～０．０１５％、
Ｓ：０．０００２～０．００１５％、
Ａｌ：０．００５～０．１５％、
Ｏ：０．０１％以下、
Ｎ：０．０１０％以下、
Ｎｂ：０．１０％以下、
Ｈ：０．０２ｐｐｍ以下を含み、
あるいはさらに、
Ｃａ：０～０．００５％、
Ｎｉ：０～２．０％、
Ｔｉ：０～０．１％、
Ｃｕ：０～１．０％、
Ｃｒ：０～１．０％、
Ｍｏ：０～０．６０％、
Ｗ：０～１．０％、
Ｖ：０～０．１０％、
Ｚｒ：０～０．０５０％、
Ｍｇ：０～０．０１％、
ＲＥＭ：０～０．０１％、
Ｂ：０～０．００２０％、
Ｔａ：０～０．２％、
Ｈｆ：０～０．２％、
Ｒｅ：０～０．００５％、
Ｓｎ：０～０．３％、
Ｓｂ：０～０．３％から選択される１種以上を含み、
残部がＦｅおよび不可避的不純物元素である、化学組成を有し、
ベイナイトおよびアスペクト比が２．０以上かつ長さが１０μｍ以上の介在物が１５個／１００ｍｍ２以下である金属組織を有し、
鋼管内面の表面から板厚中央の範囲における前記ベイナイトの最大粒径が２５μｍ以下であり、
引張強度が５２０ＭＰａ以上であって、
１ＭＰａ以上の高圧水素ガス環境において、水素誘起き裂進展下限界ＫＩＨが８０ＭＰａ・ｍ１/２以上である水素中破壊靭性に優れた高強度ラインパイプ用鋼管。
［６］　さらに、前記化学組成が、質量％で、
Ｃａ：０．０００１～０．００５％、
Ｎｉ：０．０１～２．０％、
Ｔｉ：０．００５～０．１％、
Ｃｕ：０．０１～１．０％、
Ｃｒ：０．０１～１．０％、
Ｍｏ：０．０１～０．６０％、
Ｗ：０．０１～１．０％、
Ｖ：０．０１～０．１０％、
Ｚｒ：０．０００１～０．０５０％、
Ｍｇ：０．０００１～０．０１％、
ＲＥＭ：０．０００１～０．０１％、
Ｂ：０．０００１～０．００２０％、
Ｔａ：０．０００１～０．２％、
Ｈｆ：０．０００１～０．２％、
Ｒｅ：０．０００１～０．００５％、
Ｓｎ：０．０００１～０．３％、
Ｓｂ：０．０００１～０．３％である［５］に記載の水素中破壊靭性に優れた高強度ラインパイプ用鋼管。
［７］　高強度ラインパイプ用鋼管において、
残留オーステナイトが面積分率で０～３％であり、鋼管内面の表面から板厚中央の範囲における前記ベイナイトが面積分率で９０％以上である［５］または［６］に記載の水素中破壊靭性に優れた高強度ラインパイプ用鋼管。
［８］　前記［５］または［６］に記載の成分組成を有する鋳片を１０００～１２５０℃で加熱する加熱工程と、
前記加熱工程で加熱された鋳片を、再結晶温度域での総圧下率が３５％以上５５％以下、かつ再結晶温度域での最終圧延パスの圧下率が１０％以上、かつ（再結晶温度－８０℃）以上における最終圧延パスの圧下率が１５％以上、さらに鋼板表面温度で圧延終了温度がＡｒ３変態点以上の条件で圧延する熱間圧延工程と、
該熱間圧延工程で得られた熱延鋼板を、冷却開始温度が前記熱延鋼板の表面温度でＡｒ３変態点以上、前記熱延鋼板の先端と尾端の冷却開始時間差が５０秒以内、７５０℃から５５０℃までの平均冷却速度が板厚中央温度で１５～５０℃／ｓ、冷却停止温度が２５０～６５０℃である条件で冷却する制御冷却工程と、
該制御冷却工程後、前記熱延鋼板を曲げ加工し、両端部を突合せて溶接する造管工程、前記制御冷却工程後、前記熱延鋼板を冷間ロール成形により円筒状に成形し、該円筒状の周方向両端部を突合せて電縫溶接する造管工程のうちどちらか一方の造管工程と、
を有する水素中破壊靭性に優れた高強度ラインパイプ用鋼管の製造方法。 That is, the gist of the present invention is as follows.
[1] In mass%,
C: 0.02 to 0.15%,
Si: 0.01 to 2.0%,
Mn: 0.5 to 1.5%,
P: 0.0001 to 0.015%,
S: 0.0002 to 0.0015%,
Al: 0.005 to 0.15%,
O: 0.01% or less,
N: 0.010% or less,
Nb: 0.10% or less,
H: 0.02 ppm or less,
Or even more so:
Ca: 0 to 0.005%,
Ni: 0 to 2.0%,
Ti: 0 to 0.1%,
Cu: 0 to 1.0%,
Cr: 0 to 1.0%,
Mo: 0 to 0.60%,
W: 0 to 1.0%,
V: 0 to 0.10%,
Zr: 0 to 0.050%,
Mg: 0 to 0.01%,
REM: 0 to 0.01%,
B: 0 to 0.0020%,
Ta: 0 to 0.2%,
Hf: 0 to 0.2%,
Re: 0 to 0.005%,
Sn: 0 to 0.3%,
Sb: one or more selected from 0 to 0.3%,
The balance is Fe and unavoidable impurity elements,
The metal structure has bainite and inclusions having an aspect ratio of 2.0 or more and a length of 10 μm or more at a ratio of 15 pieces/100 mm2 or less ,
The maximum grain size of the bainite in the range from the steel surface to the center of the plate thickness is 25 μm or less,
The tensile strength is 520 MPa or more,
A high-strength linepipe steel material with excellent fracture toughness in hydrogen, having a hydrogen-induced crack propagation threshold K IH of 80 MPa·m 1/2 or more in a high-pressure hydrogen gas environment of 1 MPa or more.
[2] Furthermore, the chemical composition comprises, in mass%,
Ca: 0.0001 to 0.005%,
Ni: 0.01 to 2.0%,
Ti: 0.005 to 0.1%,
Cu: 0.01 to 1.0%,
Cr: 0.01 to 1.0%,
Mo: 0.01 to 0.60%,
W: 0.01 to 1.0%,
V: 0.01 to 0.10%,
Zr: 0.0001 to 0.050%,
Mg: 0.0001 to 0.01%,
REM: 0.0001 to 0.01%,
B: 0.0001 to 0.0020%,
Ta: 0.0001 to 0.2%,
Hf: 0.0001 to 0.2%,
Re: 0.0001 to 0.005%,
Sn: 0.0001 to 0.3%,
The steel material for high-strength line pipes having excellent fracture toughness in hydrogen according to [1], wherein Sb is 0.0001 to 0.3%.
[3] A steel material for high-strength linepipes having excellent fracture toughness in hydrogen according to [1] or [2], wherein the area fraction of retained austenite is 0 to 3%, and the area fraction of the bainite in the range from the steel material surface to the center of the plate thickness is 90% or more.
[4] A heating step of heating a slab having the component composition according to [1] or [2] at 1000 to 1250 ° C.;
a hot rolling process in which the slab heated in the heating process is rolled under the conditions of a total rolling reduction in a recrystallization temperature range of 35% to 55%, a final rolling pass rolling reduction in the recrystallization temperature range of 10% or more, a final rolling pass rolling reduction at a temperature equal to or higher than (recrystallization temperature - 80 ° C.) of 15% or more, and a rolling end temperature at a steel plate surface temperature of the Ar3 transformation point or more;
A controlled cooling process in which the hot-rolled steel sheet obtained in the hot rolling process is cooled under the conditions that the cooling start temperature is the Ar3 transformation point or higher at the surface temperature of the hot-rolled steel sheet, the cooling start time difference between the front end and the tail end of the hot-rolled steel sheet is within 50 seconds, the average cooling rate from 750 ° C. to 550 ° C. is 15 to 50 ° C./s at the plate thickness center temperature, and the cooling stop temperature is 250 to 650 ° C.;
The present invention relates to a method for producing a high-strength steel material for line pipes having excellent fracture toughness in hydrogen.
[5] In mass%,
C: 0.02 to 0.15%,
Si: 0.01 to 2.0%,
Mn: 0.5 to 1.5%,
P: 0.0001 to 0.015%,
S: 0.0002 to 0.0015%,
Al: 0.005 to 0.15%,
O: 0.01% or less,
N: 0.010% or less,
Nb: 0.10% or less,
H: 0.02 ppm or less,
Or even more so:
Ca: 0 to 0.005%,
Ni: 0 to 2.0%,
Ti: 0 to 0.1%,
Cu: 0 to 1.0%,
Cr: 0 to 1.0%,
Mo: 0 to 0.60%,
W: 0 to 1.0%,
V: 0 to 0.10%,
Zr: 0 to 0.050%,
Mg: 0 to 0.01%,
REM: 0 to 0.01%,
B: 0 to 0.0020%,
Ta: 0 to 0.2%,
Hf: 0 to 0.2%,
Re: 0 to 0.005%,
Sn: 0 to 0.3%,
Sb: one or more selected from 0 to 0.3%,
The balance is Fe and unavoidable impurity elements,
The metal structure has bainite and inclusions having an aspect ratio of 2.0 or more and a length of 10 μm or more at a ratio of 15 pieces/100 mm2 or less ,
The maximum grain size of the bainite in the range from the surface of the steel pipe inner surface to the center of the plate thickness is 25 μm or less,
The tensile strength is 520 MPa or more,
A high-strength steel pipe for line pipes with excellent fracture toughness in hydrogen, having a hydrogen-induced crack propagation threshold K IH of 80 MPa·m 1/2 or more in a high-pressure hydrogen gas environment of 1 MPa or more.
[6] Furthermore, the chemical composition comprises, in mass%,
Ca: 0.0001 to 0.005%,
Ni: 0.01 to 2.0%,
Ti: 0.005 to 0.1%,
Cu: 0.01 to 1.0%,
Cr: 0.01 to 1.0%,
Mo: 0.01 to 0.60%,
W: 0.01 to 1.0%,
V: 0.01 to 0.10%,
Zr: 0.0001 to 0.050%,
Mg: 0.0001 to 0.01%,
REM: 0.0001 to 0.01%,
B: 0.0001 to 0.0020%,
Ta: 0.0001 to 0.2%,
Hf: 0.0001 to 0.2%,
Re: 0.0001 to 0.005%,
Sn: 0.0001 to 0.3%,
The steel pipe for high strength line pipe having excellent fracture toughness in hydrogen according to [5], wherein Sb is 0.0001 to 0.3%.
[7] In a high-strength steel pipe for line pipe,
7. A steel pipe for high strength line pipe having excellent fracture toughness in hydrogen according to [5] or [6], wherein the area fraction of retained austenite is 0 to 3%, and the area fraction of the bainite in the range from the surface of the inner surface of the steel pipe to the center of the plate thickness is 90% or more.
[8] A heating step of heating a slab having the component composition according to [5] or [6] at 1000 to 1250 ° C.;
a hot rolling process in which the slab heated in the heating process is rolled under the conditions of a total rolling reduction in the recrystallization temperature range of 35% to 55%, a final rolling pass rolling reduction in the recrystallization temperature range of 10% or more, a final rolling pass rolling reduction at a temperature equal to or higher than (recrystallization temperature - 80 ° C.) of 15% or more, and a rolling end temperature at the steel plate surface temperature of the Ar3 transformation point or more;
A controlled cooling process in which the hot-rolled steel sheet obtained in the hot rolling process is cooled under the following conditions: a cooling start temperature is the Ar3 transformation point or higher at the surface temperature of the hot-rolled steel sheet, a cooling start time difference between the front end and the tail end of the hot-rolled steel sheet is within 50 seconds, an average cooling rate from 750°C to 550°C is 15 to 50°C/s at the center temperature of the sheet thickness, and a cooling stop temperature is 250 to 650°C;
a pipe-making process in which, after the controlled cooling process, the hot-rolled steel sheet is bent and both ends are butted together and welded; or 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 and both circumferential ends of the cylindrical shape are butted together and electric resistance welded;
A method for manufacturing high strength steel pipe for line pipe having excellent fracture toughness in hydrogen.
　本発明によれば、高圧水素ガス環境下での水素中破壊靭性が極めて向上した鋼材を、容易にかつ簡便に製造でき、産業上格段の効果を奏する。また、本発明によれば、高圧水素ガスラインパイプ等の鋼構造物の耐水素吸収特性を顕著に向上でき、鋼構造物の安全性向上に大きく寄与するという効果もある。 According to the present invention, steel materials with extremely improved fracture toughness in hydrogen under high-pressure hydrogen gas environments can be easily and simply manufactured, which is of great industrial benefit. In addition, according to the present invention, the hydrogen absorption resistance characteristics of steel structures such as high-pressure hydrogen gas line pipes can be significantly improved, which also has the effect of greatly contributing to improving the safety of steel structures.
　次に、本発明を実施する方法について具体的に説明する。 Next, we will explain in detail how to implement the present invention.
　第１実施形態として鋼材を具体的に説明し、続いて第２実施形態として本発明の鋼管の一例であるＵＯＥ鋼管を具体的に説明し、第３実施形態として本発明の鋼管の一例である電縫鋼管を具体的に説明する。 As a first embodiment, a steel material will be specifically described, then as a second embodiment, a UOE steel pipe, which is an example of a steel pipe of the present invention, will be specifically described, and as a third embodiment, an electric resistance welded steel pipe, which is an example of a steel pipe of the present invention, will be specifically described.
　第１実施形態
　［成分組成］
　以下、本発明の鋼材における母材成分の限定理由について述べる。以下の説明において％で示す単位は、特に断らない限り全て質量％である。 First embodiment [Component composition]
The reasons for limiting the base metal components in the steel material of the present invention will be described below. In the following description, all units shown as % are mass % unless otherwise specified.
　Ｃ：０．０２～０．１５％
　Ｃは、強度の向上に有効に寄与するが、含有量が０．０２％未満では十分な強度が確保できないため、Ｃ含有量は０．０２％以上とする。好ましくは、Ｃ含有量は０．０３％以上である。より好ましくは、Ｃ含有量は０．０３５％以上である。さらに好ましくは、Ｃ含有量は０．０４％以上である。一方、０．１５％を超えると溶接性が低下する。このため、Ｃ量は０．１５％以下に限定する。好ましくは、Ｃ含有量は０．１０％以下である。また、０．０８％を超えると、制御冷却時に表層部や中心偏析部の硬さが上昇するため、耐ＳＳＣＣ性および耐ＨＩＣ性が劣化する場合がある。また、靭性も劣化する。このため、Ｃ含有量は０．０８％以下がより好ましい。さらに好ましくは、Ｃ含有量は０．０６％以下である。 C: 0.02 to 0.15%
Although C effectively contributes to improving strength, if the content is less than 0.02%, sufficient strength cannot be ensured, so the C content is set to 0.02% or more. Preferably, the C content is 0.03% or more. More preferably, the C content is 0.035% or more. Even more preferably, the C content is 0.04% or more. On the other hand, if the C content exceeds 0.15%, the weldability decreases. For this reason, the C content is limited to 0.15% or less. Preferably, the C content is 0.10% or less. Furthermore, if the C content exceeds 0.08%, the hardness of the surface layer and the central segregation increases during controlled cooling, so that the SSCC resistance and HIC resistance may deteriorate. Furthermore, the toughness also deteriorates. For this reason, the C content is more preferably 0.08% or less. Even more preferably, the C content is 0.06% or less.
　Ｓｉ：０．０１～２．０％
　Ｓｉは、脱酸のため含有するが、含有量が０．０１％未満では脱酸効果が十分でないため、Ｓｉ含有量は０．０１％以上とする。Ｓｉ含有量は、０．０２％以上が好ましい。より好ましくは、Ｓｉ含有量は０．０５％以上である。さらに好ましくは、Ｓｉ含有量は０．０８％以上である。２．０％まで上記効果が認められるため、Ｓｉ含有量は２．０％以下とする。Ｓｉ含有量は、１．８％以下が好ましく、１．５％以下がより好ましい。Ｓｉ含有量は１．０％以下がさらに好ましい。ただし、０．５％を超えると靭性や溶接性を劣化させる場合があるため、もっとも好ましくは、Ｓｉ含有量は０．５％以下とする。 Si: 0.01 to 2.0%
Si is contained for deoxidation, but if the content is less than 0.01%, the deoxidation effect is insufficient, so the Si content is set to 0.01% or more. The Si content is preferably 0.02% or more. More preferably, the Si content is 0.05% or more. Even more preferably, the Si content is 0.08% or more. Since the above effect is observed up to 2.0%, the Si content is set to 2.0% or less. The Si content is preferably 1.8% or less, more preferably 1.5% or less. Even more preferably, the Si content is 1.0% or less. However, if the Si content exceeds 0.5%, toughness and weldability may be deteriorated, so the Si content is most preferably 0.5% or less.
　Ｍｎ：０．５～１．５％
　Ｍｎは、強度、靭性の向上に有効に寄与するが、含有量が０．５％未満ではその含有効果に乏しいため、Ｍｎ含有量は０．５％以上とする。好ましくは、Ｍｎ含有量は０．６％以上であり、より好ましくは０．８％以上である。さらに好ましくは、Ｍｎ含有量は１．０％以上である。一方、１．５％を超えると制御冷却時に表層部や中心偏析部の硬さが上昇するため、耐ＳＳＣＣ性および耐ＨＩＣ性が劣化する。また、溶接性も劣化する。このため、Ｍｎ量は１．５％以下に限定する。好ましくは、Ｍｎ含有量は１．４％以下である。より好ましくは、Ｍｎ含有量は１．３％以下であり、さらに好ましくは１．２％以下である。 Mn: 0.5 to 1.5%
Mn effectively contributes to improving strength and toughness, but if the content is less than 0.5%, the effect of inclusion is poor, so the Mn content is set to 0.5% or more. Preferably, the Mn content is 0.6% or more, more preferably 0.8% or more. More preferably, the Mn content is 1.0% or more. On the other hand, if the Mn content exceeds 1.5%, the hardness of the surface layer and the central segregation increases during controlled cooling, so that the SSCC resistance and HIC resistance deteriorate. In addition, weldability also deteriorates. For this reason, the Mn content is limited to 1.5% or less. Preferably, the Mn content is 1.4% or less. More preferably, the Mn content is 1.3% or less, and even more preferably, the Mn content is 1.2% or less.
　Ｐ：０．０００１～０．０１５％
　Ｐは、不可避不純物元素であり、溶接性を劣化させるとともに、中心偏析部の硬さを上昇させることで耐ＨＩＣ性を劣化させる。０．０１５％を超えるとその傾向が顕著となるため、Ｐ含有量は０．０１５％以下に限定する。Ｐ含有量は０．０１２％以下が好ましく、０．０１０％以下がより好ましい。さらに好ましくは、Ｐ含有量は０．００８％以下である。含有量は低いほどよいが、精錬コストの観点からＰ含有量は０．０００１％以上とする。 P: 0.0001 to 0.015%
P is an inevitable impurity element that deteriorates weldability and increases the hardness of the central segregation, thereby deteriorating HIC resistance. Since this tendency becomes significant when the P content exceeds 0.015%, the P content is limited to 0.015% or less. The P content is preferably 0.012% or less, and more preferably 0.010% or less. More preferably, the P content is 0.008% or less. The lower the content, the better, but from the viewpoint of refining costs, the P content is set to 0.0001% or more.
　Ｓ：０．０００２～０．００１５％
　Ｓは、不可避不純物元素であり、鋼中においてはＭｎＳ介在物となり耐ＨＩＣ性を劣化させるため少ないことが好ましいが、０．００１５％までは許容される。このため、Ｓ含有量は０．００１５％以下とする。Ｓ含有量は０．００１０％以下が好ましく、０．０００８％以下がより好ましい。含有量は低いほどよいが、精錬コストの観点から０．０００２％以上とする。 S: 0.0002 to 0.0015%
S is an inevitable impurity element, and since it becomes MnS inclusions in steel and deteriorates HIC resistance, it is preferable that the S content is small, but up to 0.0015% is permissible. Therefore, the S content is set to 0.0015% or less. The S content is preferably 0.0010% or less, and more preferably 0.0008% or less. The lower the content, the better, but from the viewpoint of refining costs, it is set to 0.0002% or more.
　Ａｌ：０．００５～０．１５％
　Ａｌは、脱酸剤として添加するが、０．００５％未満では含有効果がないため、Ａｌ含有量は０．００５％以上である。一方、０．１５％を超えると鋼の清浄度が低下し、靱性が劣化するため、Ａｌ含有量は０．１５％以下とする。Ａｌ含有量は、０．１２％以下が好ましく、０．１０％以下がより好ましい。さらに好ましくは、Ａｌ含有量は０．０８％以下である。 Al: 0.005 to 0.15%
Al is added as a deoxidizer, but if it is less than 0.005%, there is no effect, so the Al content is 0.005% or more. On the other hand, if it exceeds 0.15%, the cleanliness of the steel decreases and the toughness deteriorates, so the Al content is set to 0.15% or less. The Al content is preferably 0.12% or less, more preferably 0.10% or less. More preferably, the Al content is 0.08% or less.
　Ｏ：０．０１％以下
　Ｏは、酸化物系介在物を生成する原因となるため少ないほど好ましい。この影響は、Ｏ含有量が０．０１％以下であれば問題とならないため、Ｏ含有量は０．０１％以下とする。Ｏ含有量は、好ましくは０．００８０％以下である。より好ましくは、Ｏ含有量は０．００３０％未満である。下限は特に限定されるものではないが、０．０００５％以上であってよい。 O: 0.01% or less O is a cause of oxide inclusions, so the less the better. This effect does not become a problem if the O content is 0.01% or less, so the O content is set to 0.01% or less. The O content is preferably 0.0080% or less. More preferably, the O content is less than 0.0030%. The lower limit is not particularly limited, but may be 0.0005% or more.
　Ｎ：０．０１０％以下
　Ｎは、強度の向上に有効に寄与するが、含有量が０．０１０％を超えると、制御冷却時に硬さが上昇するため、靭性が劣化する。このため、Ｎ含有量は０．０１０％以下とする。Ｎ含有量は０．００８％以下とすることが好ましく、Ｎ含有量は０．００６％以下とすることがより好ましく、Ｎ含有量は０．００４％以下とすることがさらに好ましい。ただし、０．００００１％未満では十分な強度が確保できないことに加え、過度の低減は製鋼上のコストを増大させる。このため、含有量は０．００００１％以上とすることが好ましい。より好ましくは、Ｎ含有量は０．００２％以上である。 N: 0.010% or less N effectively contributes to improving strength, but if the content exceeds 0.010%, the hardness increases during controlled cooling, resulting in deterioration of toughness. For this reason, the N content is set to 0.010% or less. The N content is preferably set to 0.008% or less, more preferably set to 0.006% or less, and even more preferably set to 0.004% or less. However, if the N content is less than 0.00001%, sufficient strength cannot be ensured, and excessive reduction increases the cost of steelmaking. For this reason, the content is preferably set to 0.00001% or more. More preferably, the N content is 0.002% or more.
　Ｎｂ：０．１０％以下
　Ｎｂは、鋼材の強度および靭性を高めるために有効な元素である。含有量が０．００１％未満ではその含有効果に乏しいため、０．００１％以上が好ましい。一方、０．１０％を超えると溶接部の靭性が劣化するため、Ｎｂ含有量は０．１０％以下とする。Ｎｂ含有量は０．０９５％以下とすることが好ましい。Ｎｂ含有量は０．０９０％以下とすることがより好ましく、０．０８５％以下とすることがさらに好ましい。Ｎｂ含有量は０．０８０％以下とすることがもっとも好ましい。 Nb: 0.10% or less Nb is an element effective for increasing the strength and toughness of steel. If the content is less than 0.001%, the effect of the content is poor, so 0.001% or more is preferable. On the other hand, if the content exceeds 0.10%, the toughness of the welded part deteriorates, so the Nb content is set to 0.10% or less. The Nb content is preferably set to 0.095% or less. The Nb content is more preferably set to 0.090% or less, and even more preferably set to 0.085% or less. The Nb content is most preferably set to 0.080% or less.
　Ｈ：０．０２ｐｐｍ以下
　Ｈは、製造中の種々の工程で鋼材中に導入される場合があり、導入量が多いと凝固後の割れ発生リスクが高まるとともに、ＫＩＨを著しく低下させる場合がある。これらの影響は０．０２ｐｐｍ以下であれば問題とならないため、Ｈ含有量は０．０２ｐｐｍ以下とする。Ｈ含有量は、好ましくは０．０１５ｐｐｍ以下であり、より好ましくは０．００８ｐｐｍ以下である。Ｈ含有量は、さらに好ましくは、０．００５ｐｐｍ以下、もっとも好ましくは０．００２ｐｐｍ未満である。下限は特に限定されるわけではないが、製造上のコストの理由から０．０００８ｐｐｍ以上であることが好ましい。Ｈ含有量は、０．００１ｐｐｍ以上とすることがより好ましい。なお、水素量は鋼材、鋼管、ＵＯＥ等の成形後の残存水素量である。 H: 0.02 ppm or less H may be introduced into the steel material in various processes during manufacturing, and if the amount introduced is large, the risk of cracking after solidification increases and the K IH may be significantly reduced. These effects are not a problem if the amount is 0.02 ppm or less, so the H content is set to 0.02 ppm or less. The H content is preferably 0.015 ppm or less, more preferably 0.008 ppm or less. The H content is further preferably 0.005 ppm or less, and most preferably less than 0.002 ppm. The lower limit is not particularly limited, but is preferably 0.0008 ppm or more for reasons of manufacturing costs. The H content is more preferably 0.001 ppm or more. The amount of hydrogen is the amount of hydrogen remaining after forming of steel material, steel pipe, UOE, etc.
　本開示の化学組成は、さらに、Ｃａ、Ｎｉ、Ｔｉ、Ｃｕ、Ｃｒ、Ｍｏ、Ｗ、Ｖ、Ｚｒ、Ｍｇ、ＲＥＭ、Ｂ、Ｔａ、Ｈｆ、Ｒｅ、Ｓｎ、Ｓｂのうちから選んだ１種以上を以下の範囲で任意に含有させることもできる。 The chemical composition disclosed herein may further contain one or more elements selected from Ca, Ni, Ti, Cu, Cr, Mo, W, V, Zr, Mg, REM, B, Ta, Hf, Re, Sn, and Sb in the following ranges:
　Ｃａ：０～０．００５％
　Ｃａは、硫化物系介在物の形態制御による耐ＨＩＣ性向上に有効な元素であるため、Ｃａを含有する場合には、Ｃａ含有量は０％以上であってよいが、０．０００１％未満ではその添加効果が十分でない。そのため、Ｃａを含有する場合には、Ｃａ含有量は０．０００１％以上とすることが好ましい。より好ましくは０．０００５％以上である。一方、０．００５％を超えた場合、効果が飽和するだけでなく、鋼の清浄度の低下により耐ＨＩＣ性を劣化させるので、Ｃａを含有する場合には、Ｃａ含有量は０．００５％以下に限定する。Ｃａ含有量は０．００４％以下が好ましい。Ｃａ含有量は０．００２％以下がより好ましく、０．０００８％以下がさらに好ましい。 Ca: 0 to 0.005%
Since Ca is an element effective in improving HIC resistance by controlling the morphology of sulfide-based inclusions, when Ca is contained, the Ca content may be 0% or more, but if it is less than 0.0001%, the effect of adding it is insufficient. Therefore, when Ca is contained, the Ca content is preferably 0.0001% or more. More preferably, it is 0.0005% or more. On the other hand, if it exceeds 0.005%, not only the effect is saturated, but also the HIC resistance is deteriorated due to the decrease in the cleanliness of the steel, so when Ca is contained, the Ca content is limited to 0.005% or less. The Ca content is preferably 0.004% or less. The Ca content is more preferably 0.002% or less, and even more preferably 0.0008% or less.
　Ｎｉ：０～２．０％
　Ｎｉは、靭性の改善と強度の上昇に有効な元素であり、Ｎｉを含有する場合には、Ｎｉ含有量は０％以上であってよいが、この効果を得るには０．０１％以上を含有することが好ましい。Ｎｉ含有量は０．１％以上がより好ましい。一方、コスト抑制のためには、Ｎｉを含有する場合にはＮｉ含有量は２．０％以下とする。Ｎｉ含有量は１．８％以下が好ましい。Ｎｉ含有量は１．４％以下がより好ましく、０．８％以下がさらに好ましい。 Ni: 0 to 2.0%
Ni is an element effective in improving toughness and increasing strength, and when Ni is contained, the Ni content may be 0% or more, but to obtain this effect, it is preferable to contain 0.01% or more. The Ni content is more preferably 0.1% or more. On the other hand, in order to suppress costs, when Ni is contained, the Ni content is 2.0% or less. The Ni content is preferably 1.8% or less. The Ni content is more preferably 1.4% or less, and even more preferably 0.8% or less.
　Ｔｉ：０～０．１％
　Ｔｉは、鋼材の強度上昇に寄与するため、Ｔｉを含有する場合には、Ｔｉ含有量は０％以上であってよい。前記効果を得るために、Ｔｉを含有する場合には、含有量を０．００５％以上とすることが好ましい。より好ましくは、０．００８％以上である。一方、含有量が０．１％を越えると効果が飽和し、コストアップの要因となるため、Ｔｉを含有する場合にはＴｉ含有量は０．１％以下とする。Ｔｉ含有量は０．０８％以下とすることが好ましく、０．０６％以下とすることがより好ましい。コスト抑制のためには、Ｔｉ含有量は０．０５％以下とすることがさらに好ましい。Ｔｉ含有量は０．０４％以下とすることがもっとも好ましい。 Ti: 0 to 0.1%
Since Ti contributes to increasing the strength of the steel material, when Ti is contained, the Ti content may be 0% or more. In order to obtain the above effect, when Ti is contained, the content is preferably 0.005% or more. More preferably, it is 0.008% or more. On the other hand, when the content exceeds 0.1%, the effect is saturated and becomes a factor of increasing costs, so when Ti is contained, the Ti content is 0.1% or less. The Ti content is preferably 0.08% or less, and more preferably 0.06% or less. In order to suppress costs, the Ti content is further preferably 0.05% or less. The Ti content is most preferably 0.04% or less.
　Ｃｕ：０～１．０％
　Ｃｕは、靭性の改善と強度の上昇に有効な元素であり、Ｃｕを含有する場合には、Ｃｕ含有量は０％以上であってよいが、この効果を得るには０．０１％以上を含有することが好ましい。０．０５％以上とすることがより好ましい。一方、含有量が多すぎると溶接性が劣化するため、Ｃｕを含有する場合は、Ｃｕ含有量は１．０％以下とする。Ｃｕ含有量は０．９５％以下が好ましく、Ｃｕ含有量は０．９％以下がより好ましい。さらに好ましくは、Ｃｕ含有量は０．８５％以下である。もっとも好ましくは、Ｃｕ含有量は０．５％以下である。 Cu: 0 to 1.0%
Cu is an element effective in improving toughness and increasing strength, and when Cu is contained, the Cu content may be 0% or more, but to obtain this effect, it is preferable to contain 0.01% or more. It is more preferable to make it 0.05% or more. On the other hand, if the content is too high, weldability deteriorates, so when Cu is contained, the Cu content is 1.0% or less. The Cu content is preferably 0.95% or less, and more preferably 0.9% or less. More preferably, the Cu content is 0.85% or less. Most preferably, the Cu content is 0.5% or less.
　Ｃｒ：０～１．０％
　Ｃｒは、Ｍｎと同様、低Ｃでも十分な強度を得るために有効な元素であり、Ｃｒを含有する場合には、Ｃｒ含有量は０％以上であってよいが、この効果を得るには０．０１％以上を含有することが好ましい。０．０５％以上とすることがより好ましい。一方、含有量が多すぎると、焼入れ性が過剰になるため、耐ＳＳＣＣ性が劣化する。また、溶接性も劣化する。このため、Ｃｒを含有する場合は１．０％以下とする。Ｃｒ含有量は０．９５％以下が好ましい。Ｃｒ含有量は０．９％以下がより好ましく、０．８５％以下がさらに好ましい。 Cr: 0 to 1.0%
Like Mn, Cr is an effective element for obtaining sufficient strength even with low C. When Cr is contained, the Cr content may be 0% or more, but to obtain this effect, it is preferable to contain 0.01% or more. It is more preferable to make it 0.05% or more. On the other hand, if the content is too high, the hardenability becomes excessive, so that the SSCC resistance deteriorates. In addition, the weldability also deteriorates. Therefore, when Cr is contained, it is 1.0% or less. The Cr content is preferably 0.95% or less. The Cr content is more preferably 0.9% or less, and even more preferably 0.85% or less.
　Ｍｏ：０～０．６０％
　Ｍｏは、靭性の改善と強度の上昇に有効な元素であり、耐ＳＳＣＣ性、耐ＨＩＣ性の向上に有効な元素である。Ｍｏを含有する場合には、Ｍｏ含有量は０％以上であってよいが、この効果を得るには０．０１％以上を含有することが好ましい。０．１０％以上を含有することがより好ましい。一方で、含有量が多すぎると、焼入れ性が過剰になるため、耐ＳＳＣＣ性が劣化する。また、溶接性も劣化する。このため、Ｍｏを含有する場合には、Ｍｏ含有量は０．６０％以下とする。Ｍｏ含有量は、好ましくは０．５０％以下とする。より好ましくは０．４０％以下とする。さらに好ましくは０．３５％以下とする。 Mo: 0 to 0.60%
Mo is an element effective in improving toughness and increasing strength, and is an element effective in improving SSCC resistance and HIC resistance. When Mo is contained, the Mo content may be 0% or more, but to obtain this effect, it is preferable to contain 0.01% or more. It is more preferable to contain 0.10% or more. On the other hand, if the content is too high, the hardenability becomes excessive, so that the SSCC resistance deteriorates. In addition, the weldability also deteriorates. Therefore, when Mo is contained, the Mo content is 0.60% or less. The Mo content is preferably 0.50% or less. More preferably, it is 0.40% or less. Further preferably, it is 0.35% or less.
　Ｗ：０～１．０％
　Ｗは、鋼材の強度上昇に寄与する。Ｗを含有する場合には、Ｗ含有量は０％以上であってよいが、前記効果を得るために、Ｗを含有する場合には、含有量を０．０１％以上とすることが好ましい。一方、Ｗ含有量が１．０％を越えると効果が飽和し、コストアップの要因となるため、Ｗを含有する場合には、Ｗ含有量は１．０％以下とする。Ｗ含有量は０．９％以下とすることが好ましく、０．８％以下とすることがより好ましい。コスト抑制のためには、０．５％以下とすることがさらに好ましい。 W: 0 to 1.0%
W contributes to increasing the strength of the steel material. When W is contained, the W content may be 0% or more, but in order to obtain the above effect, when W is contained, the content is preferably 0.01% or more. On the other hand, when the W content exceeds 1.0%, the effect is saturated and becomes a factor of increasing costs, so when W is contained, the W content is 1.0% or less. The W content is preferably 0.9% or less, and more preferably 0.8% or less. In order to suppress costs, it is even more preferable to make it 0.5% or less.
　Ｖ：０～０．１０％、Ｚｒ：０～０．０５０％、Ｍｇ：０～０．０１％およびＲＥＭ：０～０．０１％
　Ｖは、鋼材の強度および靭性を高めるために任意に含有することができる元素である。Ｖを含有する場合には、Ｖ含有量は０％以上であってよいが、含有量が０．０１％未満ではその含有効果に乏しいため、Ｖ含有量は０．０１％以上とすることが好ましい。Ｖ含有量は０．０３％以上とすることがより好ましい。一方、０．１０％を超えると溶接部の靭性が劣化するので、含有する場合は０．１０％以下とするのが好ましい。Ｖ含有量が０．０９％以下とすることが好ましい。Ｖ含有量が０．０７％以下とすることがより好ましく、０．０６％以下とすることがさらに好ましい。 V: 0 to 0.10%, Zr: 0 to 0.050%, Mg: 0 to 0.01%, and REM: 0 to 0.01%.
V is an element that can be optionally contained to increase the strength and toughness of the steel material. When V is contained, the V content may be 0% or more, but if the content is less than 0.01%, the effect of the inclusion is poor, so the V content is preferably 0.01% or more. The V content is more preferably 0.03% or more. On the other hand, if it exceeds 0.10%, the toughness of the weld deteriorates, so if it is contained, it is preferably 0.10% or less. The V content is preferably 0.09% or less. The V content is more preferably 0.07% or less, and even more preferably 0.06% or less.
　Ｚｒ、ＭｇおよびＲＥＭは、結晶粒微細化を通じて靭性を高めたり、介在物性状のコントロールを通して耐割れ性を高めたりするために任意に添加することができる元素である。これらの元素を含有する場合には、その含有量は０％以上であってよいが、いずれも、含有量が０．０００１％未満ではその含有効果に乏しいため、含有量は０．０００１％以上とすることが好ましい。より好ましくは０．０００５％以上である。すなわち、Ｚｒ含有量は０．０００１％以上とすることが好ましい。Ｚｒ含有量は０．０００５％以上とすることがより好ましい。また、ＲＥＭ含有量は０．０００１％以上とすることが好ましい。ＲＥＭ含有量は０．０００５％以上とすることがより好ましい。Ｍｇ含有量は０．０００１％以上とすることが好ましい。Ｍｇ含有量は０．０００５％以上とすることがより好ましい。 Zr, Mg and REM are elements that can be added at will to improve toughness through grain refinement and crack resistance through control of inclusion properties. When these elements are contained, their content may be 0% or more, but since the effect of containing them is poor when the content is less than 0.0001%, the content is preferably 0.0001% or more. More preferably, it is 0.0005% or more. In other words, the Zr content is preferably 0.0001% or more. The Zr content is more preferably 0.0005% or more. The REM content is preferably 0.0001% or more. The REM content is more preferably 0.0005% or more. The Mg content is preferably 0.0001% or more. The Mg content is more preferably 0.0005% or more.
　一方、Ｚｒ含有量は０．０５０％を超える、またＭｇおよびＲＥＭ含有量は、０．０１％を超えるとその効果が飽和するので、含有する場合は０．０５０％以下、ＭｇおよびＲＥＭ含有量は０．０１％以下とする。すなわち、含有する場合には、Ｚｒ含有量は０．０５０％以下とする。Ｚｒ含有量は０．０４０％以下とすることが好ましい。Ｚｒ含有量は０．０２０％以下とすることがより好ましい。また、含有する場合には、ＲＥＭ含有量は０．０１％以下とする。ＲＥＭ含有量は０．００９％以下とすることが好ましい。ＲＥＭ含有量は０．００８％以下とすることがより好ましい。また、含有する場合には、Ｍｇ含有量は０．０１％以下とする。Ｍｇ含有量は０．００９％以下とすることが好ましい。Ｍｇ含有量は０．００８％以下とすることがより好ましい。 On the other hand, if the Zr content exceeds 0.050%, and if the Mg and REM contents exceed 0.01%, the effect saturates, so if they are contained, they should be 0.050% or less, and the Mg and REM contents should be 0.01% or less. That is, if they are contained, the Zr content should be 0.050% or less. It is preferable that the Zr content be 0.040% or less. It is more preferable that the Zr content be 0.020% or less. If they are contained, the REM content should be 0.01% or less. It is preferable that the REM content be 0.009% or less. It is more preferable that the REM content be 0.008% or less. If they are contained, the Mg content should be 0.01% or less. It is preferable that the Mg content be 0.009% or less. It is more preferable that the Mg content be 0.008% or less.
　Ｂ：０～０．００２０％
　Ｂは、焼き入れ性を向上させる元素であり、鋼材の強度上昇に寄与するとともに、旧オーステナイト粒の粗大化を抑制し、素材の各種特性を向上させる。Ｂを含有する場合には、Ｂ含有量は０％以上であってよいが、前記効果を得るために、含有量を０．０００１％以上とすることが好ましい。より好ましくは、０．０００８％以上である。一方、Ｂ含有量が０．００２０％を越えると効果が飽和し、コストアップの要因となるため、Ｂを含有する場合には、Ｂ含有量は０．００２０％以下とする。Ｂ含有量は０．００１４％以下とすることが好ましい。Ｂ含有量は０．００１２％以下とすることがより好ましい。コスト抑制のためには、０．００１０％以下とすることがさらに好ましい。 B: 0 to 0.0020%
B is an element that improves hardenability, contributes to increasing the strength of the steel material, inhibits the coarsening of prior austenite grains, and improves various properties of the material. When B is contained, the B content may be 0% or more, but in order to obtain the above effect, the content is preferably 0.0001% or more. More preferably, it is 0.0008% or more. On the other hand, when the B content exceeds 0.0020%, the effect is saturated and becomes a factor of increasing costs, so when B is contained, the B content is 0.0020% or less. The B content is preferably 0.0014% or less. The B content is more preferably 0.0012% or less. In order to suppress costs, it is even more preferably 0.0010% or less.
　Ｔａ：０～０．２％
　Ｔａは、炭化物、窒化物を形成し強度の向上に寄与する元素である。Ｔａを含有する場合には、Ｔａ含有量は０％以上であってよいが、前記効果を得るために、Ｔａ含有量は０．０００１％以上を含有することが好ましい。より好ましくは、Ｔａ含有量は、０．０００８％以上である。一方、含有量が０．２％を超えると靭性の低下を招くことがあるため、Ｔａを含有する場合には、Ｔａ含有量は０．２％以下とする。Ｔａは０．１６％以下とすることが好ましい。Ｔａは０．１２％以下とすることがより好ましく、０．１０％以下とすることがさらに好ましい。 Ta: 0 to 0.2%
Ta is an element that forms carbides and nitrides and contributes to improving strength. When Ta is contained, the Ta content may be 0% or more, but in order to obtain the above effect, it is preferable that the Ta content is 0.0001% or more. More preferably, the Ta content is 0.0008% or more. On the other hand, if the content exceeds 0.2%, it may cause a decrease in toughness, so when Ta is contained, the Ta content is 0.2% or less. It is preferable that Ta is 0.16% or less. It is more preferable that Ta is 0.12% or less, and even more preferable that Ta is 0.10% or less.
　Ｈｆ：０～０．２％、Ｒｅ：０～０．００５％
　これらの元素は、鋼材の強度上昇に寄与する。前記効果を得るために、含有する場合には、これらの元素の含有量を０．０００１％以上とすることが好ましい。好ましくは、０．００１０％以上である。すなわち、含有する場合には、Ｈｆ含有量は、０．０００１％以上が好ましい。Ｈｆ含有量は０．００１０％以上がより好ましい。含有する場合には、Ｒｅ含有量は０．０００１％以上が好ましい。Ｒｅ含有量は０．００１％以上が好ましい。一方、これらの元素を含有する場合には、含有量がＨｆは０．２％を超える、Ｒｅは０．００５％を越えると、酸化物が増加し、凝集すると耐水素特性を損なうため、Ｈｆは０．２％以下、Ｒｅ含有量は０．００５％以下とする。すなわち、含有する場合には、Ｈｆ含有量は０．２％以下とする。Ｈｆ含有量は、０．１８％以下とすることが好ましく、０．１２％以下とすることがより好ましい。含有する場合には、Ｒｅ含有量は０．００５％以下とする。Ｒｅ含有量は、０．００４％以下とすることが好ましく、０．００３％以下とすることがより好ましい。 Hf: 0 to 0.2%, Re: 0 to 0.005%
These elements contribute to increasing the strength of the steel material. In order to obtain the above-mentioned effect, when these elements are contained, the contents of these elements are preferably 0.0001% or more. Preferably, they are 0.0010% or more. That is, when these elements are contained, the Hf content is preferably 0.0001% or more. The Hf content is more preferably 0.0010% or more. When these elements are contained, the Re content is preferably 0.0001% or more. The Re content is preferably 0.001% or more. On the other hand, when these elements are contained, if the content of Hf exceeds 0.2% and the content of Re exceeds 0.005%, oxides increase and, if they aggregate, hydrogen resistance properties are impaired, so Hf is 0.2% or less and the Re content is 0.005% or less. That is, when these elements are contained, the Hf content is 0.2% or less. The Hf content is preferably 0.18% or less, and more preferably 0.12% or less. When these elements are contained, the Re content is 0.005% or less. The Re content is preferably 0.004% or less, and more preferably 0.003% or less.
　Ｓｎ：０～０．３％、Ｓｂ：０～０．３％、
　これらの元素は、鋼材の強度上昇と焼入れ性向上に寄与する。Ｓｎ、Ｓｂを含有する場合には、Ｓｎ、Ｓｂ含有量は０％以上であってよいが、前記効果を得るために、含有量をそれぞれ０．０００１％以上とすることが好ましい。好ましくは、０．００１％以上である。すなわち、Ｓｎ含有量は、含有する場合には０％以上であってよいが、Ｓｎ含有量は０．０００１％以上とすることが好ましい。Ｓｎ含有量は、０．００１％以上とすることがより好ましい。Ｓｂ含有量は、含有する場合には０％以上であってよいが、Ｓｂ含有量は０．０００１％以上とすることが好ましい。Ｓｂ含有量は、０．００１％以上とすることがより好ましい。
一方、それぞれ含有量が０．３％を越えると効果が飽和し、コストアップの要因となるため、Ｓｎ、Ｓｂを含有する場合には、Ｓｎ、Ｓｂ含有量は０．３％以下とする。コスト抑制のためには、０．０１％以下とすることが好ましい。すなわち、Ｓｎを含有する場合には、Ｓｎ含有量は０．３％以下とする。Ｓｎ含有量は０．２％以下とすることが好ましい。Ｓｎ含有量は０．１％以下とすることがより好ましい。Ｓｎ含有量は０．０１％以下とすることがさらに好ましい。Ｓｂを含有する場合には、Ｓｂ含有量は０．３％以下とする。Ｓｂ含有量は０．２％以下とすることが好ましい。Ｓｂ含有量は０．１％以下とすることがより好ましい。Ｓｂ含有量は０．０１％以下とすることがさらに好ましい。 Sn: 0 to 0.3%, Sb: 0 to 0.3%,
These elements contribute to increasing the strength and hardenability of the steel material. When Sn and Sb are contained, the Sn and Sb contents may be 0% or more, but in order to obtain the above effects, it is preferable that the contents are each 0.0001% or more. Preferably, they are 0.001% or more. That is, when Sn is contained, the Sn content may be 0% or more, but it is preferable that the Sn content is 0.0001% or more. It is more preferable that the Sn content is 0.001% or more. When Sb is contained, the Sb content may be 0% or more, but it is preferable that the Sb content is 0.0001% or more. It is more preferable that the Sb content is 0.001% or more.
On the other hand, when the content exceeds 0.3%, the effect is saturated and becomes a factor of cost increase, so when Sn and Sb are contained, the Sn and Sb contents are 0.3% or less. In order to suppress costs, it is preferable to make it 0.01% or less. That is, when Sn is contained, the Sn content is 0.3% or less. It is preferable to make the Sn content 0.2% or less. It is more preferable to make the Sn content 0.1% or less. It is even more preferable to make the Sn content 0.01% or less. When Sb is contained, the Sb content is 0.3% or less. It is preferable to make the Sb content 0.2% or less. It is more preferable to make the Sb content 0.1% or less. It is even more preferable to make the Sb content 0.01% or less.
　鋼材の成分組成において、上述した成分（元素）以外の残部は、Ｆｅおよび不可避的不純物元素からなる。 In the composition of the steel material, the remainder other than the above-mentioned components (elements) consists of Fe and unavoidable impurity elements.
　以下、本発明の鋼材の金属組織について述べる。 The metal structure of the steel material of the present invention is described below.
　金属組織
　アスペクト比が２．０以上かつ長さが１０μｍ以上の介在物が１５個／１００ｍｍ２以下
　材料中の介在物として、例えば伸長したＭｎＳ、セメンタイトが挙げられる。これらは、水素集積源として作用し、耐ＨＩＣ性の著しい低下を招くとともに、水素誘起き裂進展下限界ＫＩＨ低下の原因となる。このため、アスペクト比が２．０以上かつ長さが１０μｍ以上の介在物が１５個／１００ｍｍ２以下とする。上記介在物の個数密度は、好ましくは１０個／１００ｍｍ２以下である。下限は特に限定されるものではなく、０個／１００ｍｍ２であってもよい。 Metal structure: 15 inclusions/ 100 mm2 or less with an aspect ratio of 2.0 or more and a length of 10 μm or more Examples of inclusions in the material include elongated MnS and cementite. These act as hydrogen accumulation sources, leading to a significant decrease in HIC resistance and causing a decrease in the hydrogen-induced crack propagation limit K IH . For this reason, the number of inclusions with an aspect ratio of 2.0 or more and a length of 10 μm or more is set to 15 inclusions/100 mm2 or less. The number density of the inclusions is preferably 10 inclusions/100 mm2 or less. The lower limit is not particularly limited, and may be 0 inclusions/100 mm2 .
　残留オーステナイトが０～３％（好適）
　残留オーステナイトが鋼材組織中に残存することにより、水素トラップサイトとして作用し、鋼中の水素量が増大し、水素脆化感受性を増大させる場合がある。さらに、鋼材や鋼管を鋼構造物として使用する場合、使用中の応力負荷により、残留オーステナイトがマルテンサイトに変態した場合、マルテンサイトは非常に硬質であり、ＨＩＣの発生源あるいは伝播経路となり、ＫＩＨを著しく低下させる可能性がある。本発明においては、残留オーステナイトを３％以下とすることで、ＫＩＨを向上させた。このため、残留オーステナイトは３％以下とすることが好ましい。残留オーステナイトは２％以下とすることがより好ましい。さらに好ましくは１％以下である。残留オーステナイトは０％であってもよい。 Retained austenite is 0-3% (preferable)
When retained austenite remains in the steel structure, it acts as a hydrogen trap site, increasing the amount of hydrogen in the steel and increasing the hydrogen embrittlement susceptibility. Furthermore, when a steel material or steel pipe is used as a steel structure, if the retained austenite transforms into martensite due to stress load during use, the martensite is very hard and may become a source or propagation path of HIC, significantly decreasing K IH . In the present invention, the retained austenite is set to 3% or less to improve K IH . For this reason, the retained austenite is preferably set to 3% or less. The retained austenite is more preferably set to 2% or less. More preferably, it is set to 1% or less. The retained austenite may be 0%.
　鋼材表面（鋼管の場合は鋼管内面の表面）から板厚中央の範囲におけるベイナイトが面積分率で９０％以上（好適）
　ラインパイプに好適な材料として、引張強さが５２０ＭＰａ以上の高強度化を図るために、鋼材の鋼組織はベイナイト組織とする必要がある。ここで、ベイナイト組織は、変態強化に寄与する加速冷却時あるいは加速冷却後に変態するベイニティックフェライトまたはグラニュラーベイナイトを含み、かつ、焼き戻しベイナイトを含むものとする。ベイナイト組織中に、フェライト、マルテンサイト、パーライト、島状マルテンサイトおよび残留オーステナイトなどの異種組織が混在すると、強度の低下や通常（大気環境下での）の靭性、およびＫＩＨの劣化が生じる。
また、異なる硬度の鋼組織の存在は、使用中の応力負荷時に鋼材内に応力分布を生じる原因となり、応力誘起拡散による水素集積源として作用し、耐ＨＩＣ特性を低下させる。このことから、ベイナイトを面積分率で９０％以上とすることが好ましい。ベイナイトは面積分率で９２％以上とすることがより好ましく、９５％以上とすることがさらに好ましい。上限については特に限定されるものではないが、１００％であってもよい。 The area fraction of bainite in the range from the steel surface (in the case of steel pipes, the surface of the steel pipe inner surface) to the center of the plate thickness is 90% or more (preferred).
In order to achieve a high strength of tensile strength of 520 MPa or more as a material suitable for line pipes, the steel material must have a bainite structure. Here, the bainite structure includes bainitic ferrite or granular bainite that transforms during or after accelerated cooling, which contributes to transformation strengthening, and also includes tempered bainite. If the bainite structure includes heterogeneous structures such as ferrite, martensite, pearlite, island martensite, and retained austenite, the strength decreases, and the normal (atmospheric) toughness and K IH deteriorate.
Furthermore, the presence of steel structures with different hardnesses causes stress distribution in the steel material when stress is applied during use, and acts as a hydrogen accumulation source due to stress-induced diffusion, thereby degrading HIC resistance. For this reason, it is preferable that bainite is 90% or more in area fraction. It is more preferable that bainite is 92% or more in area fraction, and even more preferable that bainite is 95% or more in area fraction. There is no particular upper limit, but it may be 100%.
　鋼材表面（鋼管の場合は鋼管内面の表面）から板厚中央の範囲における最大粒径が２５μｍ以下
　平均結晶粒径を微細化することで、靭性は向上するが、Ａｒ３点以上で冷却を開始する場合においては、平均結晶粒径の微細化に限界がある。本開示においては、粗大な結晶粒の形成を抑制することが肝要である。最大粒径が大きい結晶粒を含む場合、材料内部に不均一なひずみの発生を誘起し、水素の集積を助長するため、水素ガス環境における破壊靭性が劣化する。特に鋼材内面の表面から板厚中央における最大粒径が２５μｍ超である結晶粒は、粒の周囲にひずみを蓄積しやすく、容易に水素き裂の起点、および進展経路となるので、ＫＩＨが著しく劣化する。よって、鋼材内表面から板厚中央における最大粒径を２５μｍ以下とする必要がある。鋼材内表面から板厚中央における最大粒径は２４μｍ以下とすることが好ましく、２２μｍ以下とすることがより好ましく、２０μｍ以下とすることがさらに好ましい。下限は特に限定されるものではないが、最大粒径は４μｍ以上とすることが好ましい。結晶粒径の測定範囲は、１ｍｍ×１ｍｍとし、結晶粒径の定義はＡｒｅａ粒径（方位差が１５°以上の境界を粒界と定義したときの加重平均）とした。 The maximum grain size in the range from the steel surface (the surface of the steel pipe inner surface in the case of a steel pipe) to the center of the plate thickness is 25 μm or less. By refining the average grain size, toughness is improved, but when cooling is started at the Ar 3 point or higher, there is a limit to refining the average grain size. In the present disclosure, it is essential to suppress the formation of coarse grains. When grains with a large maximum grain size are included, the occurrence of non-uniform strain is induced inside the material, promoting the accumulation of hydrogen, and therefore the fracture toughness in a hydrogen gas environment is deteriorated. In particular, grains with a maximum grain size of more than 25 μm from the surface of the inner steel material to the center of the plate thickness tend to accumulate strain around the grains, which easily become the starting point and propagation path of hydrogen cracks, and therefore the K IH is significantly deteriorated. Therefore, it is necessary to make the maximum grain size from the inner surface of the steel material to the center of the plate thickness 25 μm or less. The maximum grain size from the inner surface of the steel material to the center of the plate thickness is preferably 24 μm or less, more preferably 22 μm or less, and even more preferably 20 μm or less. Although the lower limit is not particularly limited, the maximum grain size is preferably 4 μm or more. The measurement range of the crystal grain size is 1 mm × 1 mm, and the crystal grain size is defined as the area grain size (weighted average when the boundary with an orientation difference of 15° or more is defined as the grain boundary).
　１ＭＰａ以上の高圧水素ガス環境における水素誘起き裂進展下限界ＫＩＨが８０ＭＰａ・ｍ１／２以上
　本開示の高強度鋼材は、水素を含有する環境における鋼構造物の安全な運用のために、１ＭＰａ以上の高圧水素ガス環境において、鋼材の水素誘起き裂進展下限界ＫＩＨが８０ＭＰａ・ｍ１／２以上を有するものとする。上限は特に限定されるものではないが、鋼材の水素誘起き裂進展下限界ＫＩＨが１２０ＭＰａ・ｍ１／２以下を有することが好ましく、１００Ｐａ・ｍ１／２以下とすることがより好ましい。なお、水素誘起き裂進展下限界ＫＩＨは、１ＭＰａ以上の高圧水素ガス中においてＡＳＴＭ　Ｅ３９９、　ＡＳＴＭ　Ｅ１８２０に準拠し求めた平面ひずみ破壊靭性ＫＩＣまたはその暫定値を指し、あるいはＡＳＴＭ　Ｅ１６８１に準拠し求めたき裂進展下限界値またはその暫定値とする。 The hydrogen-induced crack propagation threshold K IH in a high-pressure hydrogen gas environment of 1 MPa or more is 80 MPa·m 1/2 or more. In order to safely operate steel structures in a hydrogen-containing environment, the high-strength steel material of the present disclosure has a hydrogen-induced crack propagation threshold K IH of 80 MPa·m 1/2 or more in a high-pressure hydrogen gas environment of 1 MPa or more. Although the upper limit is not particularly limited, it is preferable that the hydrogen-induced crack propagation threshold K IH of the steel material is 120 MPa·m 1/2 or less, and more preferably 100 Pa·m 1/2 or less. The hydrogen-induced crack propagation threshold K IH refers to the plane strain fracture toughness K IC or its provisional value obtained in a high-pressure hydrogen gas of 1 MPa or more in accordance with ASTM E399 or ASTM E1820, or the crack propagation threshold or its provisional value obtained in accordance with ASTM E1681.
　また、鋼材の板厚は特に限定されるものではないが、板厚は５ｍｍ以上とすることが好ましい。板厚は３０ｍｍ以下とすることが好ましい。 The thickness of the steel plate is not particularly limited, but it is preferable that the thickness be 5 mm or more. It is preferable that the thickness be 30 mm or less.
　本発明は、上述した化学成分と、金属組織を有することで，優れた高圧水素ガス中における水素誘起き裂下限界ＫＩＨを得ることができ、水素ラインパイプへの適用が可能となる。 The present invention, by having the above-mentioned chemical composition and metal structure, can obtain an excellent hydrogen-induced cracking threshold K IH in high-pressure hydrogen gas, and can be applied to hydrogen line pipes.
　さらに、本発明に係るラインパイプ用高強度鋼材は、下記に示す製造条件を限定することにより得ることができ、製造方法および条件を具体的に説明する。 Furthermore, the high-strength steel for line pipes according to the present invention can be obtained by limiting the manufacturing conditions shown below, and the manufacturing method and conditions are specifically explained below.
　溶鋼工程
　［溶鋼の平均冷却速度：５０℃／分以上（好適条件）］
　介在物の低減には、ＳあるいはＯの含有量を低減することも有効である。溶鋼の冷却過程で、本発明で限定している介在物は凝集するため、溶鋼の平均冷却速度を速くすることも有効である。このため、１５００℃から１０００℃までの温度域における平均冷却速度を５０℃／分以上とすることが好ましい。平均冷却速度は６０℃／分以上とすることがより好ましく、７０℃／分以上とすることがさらに好ましい。上限は特に限定されるものではないが、９０℃／分以下にすることが好ましい。 Molten steel process [Average cooling rate of molten steel: 50° C./min or more (preferred conditions)]
In order to reduce the number of inclusions, it is also effective to reduce the content of S or O. Since the inclusions defined in the present invention aggregate during the cooling process of the molten steel, it is also effective to increase the average cooling rate of the molten steel. For this reason, it is preferable to set the average cooling rate in the temperature range from 1500°C to 1000°C to 50°C/min or more. It is more preferable to set the average cooling rate to 60°C/min or more, and even more preferable to set the average cooling rate to 70°C/min or more. There is no particular upper limit, but it is preferable to set the rate to 90°C/min or less.
　加熱工程
　［鋳片の加熱温度：１０００～１２５０℃］
　ビレットやスラブ等の鋳片加熱温度は、１０００℃未満ではミクロ偏析しているＣやＰ、Ｓ等の不純物元素の拡散が不十分で均質な材質が得られないため、介在物数の増大，および不均一析出を引き起こし、靭性を低下させる。そのため、鋳片の加熱温度は１０００℃以上とする。鋳片の加熱温度は１０５０℃以上とすることが好ましく、１１００℃以上とすることがより好ましい。一方、１２５０℃を超えると、結晶粒が粗大化しすぎ靱性が劣化する。従って、鋳片の加熱温度は１２５０℃以下とする。鋳片の加熱温度は１２００℃以下とすることが好ましく、１１５０℃以下とすることがより好ましい。 Heating process [heating temperature of cast piece: 1000 to 1250°C]
If the heating temperature of the billet or slab is less than 1000°C, the diffusion of micro-segregated impurity elements such as C, P, and S is insufficient, and a homogeneous material cannot be obtained, which causes an increase in the number of inclusions and non-uniform precipitation, thereby reducing toughness. Therefore, the heating temperature of the billet is set to 1000°C or higher. The heating temperature of the billet is preferably set to 1050°C or higher, and more preferably set to 1100°C or higher. On the other hand, if the heating temperature exceeds 1250°C, the crystal grains become too coarse, and the toughness is deteriorated. Therefore, the heating temperature of the billet is set to 1250°C or lower. The heating temperature of the billet is preferably set to 1200°C or lower, and more preferably set to 1150°C or lower.
　圧延工程
　［鋳片加熱後の再結晶温度域での総圧下率：３５％以上５５％以下］
　ベイナイトの最大粒径を小さくするためには、鋳片加熱後の再結晶温度域での熱間圧延で、結晶粒の再結晶を促進し、粗大粒の形成を抑制する必要がある。再結晶温度域での総圧下率が３５％未満の場合、再結晶が不十分であるため、粗大粒が残存する。よって、再結晶温度域での総圧下率は３５％以上とする。好ましくは３８％以上とする。再結晶温度域での総圧下率は、より好ましくは４０％以上であり、さらに好ましくは４３％以上とする。一方、再結晶温度域での総圧下率が５５％を超えると、結晶粒の粗大化は抑制できるが、未再結晶域での圧下が不足するため、最終製品における結晶粒の微細化ができない。よって、再結晶温度域での総圧下率は５５％以下とする。好ましくは５２％以下とする。再結晶温度域での総圧下率は、より好ましくは５０％以下であり、さらに好ましくは４８％以下とする。ここで、再結晶の下限温度Ｔｎｒは、例えば、鋼の成分から以下の式で求めることができる。なお、鋼板の表面温度は放射温度計等で測定することができる。再結晶温度域での総圧下率とは、下記式により求めた再結晶の下限温度Ｔｎｒ以上での合計圧下率をさす。
Ｔｎｒ（℃）＝１７４×ｌｏｇ[％Ｎｂ][％Ｃ+１２／１４％Ｎ]＋１４４４
ただし、[％Ｘ]はＸ元素の鋼中含有量（質量％）を示す。 Rolling process [Total rolling reduction in the recrystallization temperature range after heating of the slab: 35% to 55%]
In order to reduce the maximum grain size of bainite, it is necessary to promote the recrystallization of crystal grains and suppress the formation of coarse grains by hot rolling in the recrystallization temperature range after heating of the slab. If the total reduction in the recrystallization temperature range is less than 35%, the recrystallization is insufficient and coarse grains remain. Therefore, the total reduction in the recrystallization temperature range is set to 35% or more. Preferably, it is set to 38% or more. The total reduction in the recrystallization temperature range is more preferably set to 40% or more, and even more preferably set to 43% or more. On the other hand, if the total reduction in the recrystallization temperature range exceeds 55%, the coarsening of crystal grains can be suppressed, but the reduction in the non-recrystallized range is insufficient, so the crystal grains in the final product cannot be refined. Therefore, the total reduction in the recrystallization temperature range is set to 55% or less. Preferably, it is set to 52% or less. The total reduction in the recrystallization temperature range is more preferably set to 50% or less, and even more preferably set to 48% or less. Here, the lower limit temperature Tnr of recrystallization can be calculated, for example, from the components of the steel by the following formula. The surface temperature of the steel sheet can be measured by a radiation thermometer, etc. The total reduction in the recrystallization temperature range refers to the total reduction at or above the lower limit temperature Tnr of recrystallization calculated by the following formula.
Tnr(℃)=174×log[%Nb][%C+12/14%N]+1444
Here, [%X] indicates the content (mass%) of the X element in the steel.
　［再結晶温度域での最終圧延パスの圧下率：１０％以上］
　上述の再結晶温度域での総圧下率を３５％以上５５％以下とするのに加えて、再結晶温度域での最終圧延パスの圧下率を十分に確保し、再結晶を十分に促進させることで、粗大粒が存在しない均一粒の状態で部分再結晶域圧延を開始する必要がある。再結晶温度域での最終圧延パスの圧下率が１０％未満の場合、再結晶が不十分であるため、粗圧延後仕上げ圧延開始までの保持時間の間に粗大粒に成長する。よって、再結晶温度域での最終圧延パスの圧下率は１０％以上とする。再結晶温度域での最終圧延パスの圧下率は好ましくは１１％以上とする。再結晶温度域での最終圧延パスの圧下率は、より好ましくは１３％以上であり、さらに好ましくは１５％以上である。再結晶温度域での最終圧延パスの圧下率の上限は特に限定されなく、高いほど好ましいが、７０％を超えると生産性が著しく低下するため、７０％以下とすることが好ましい。 [Reduction rate of final rolling pass in recrystallization temperature range: 10% or more]
In addition to the total reduction in the recrystallization temperature range being 35% or more and 55% or less, it is necessary to ensure a sufficient reduction in the final rolling pass in the recrystallization temperature range and sufficiently promote recrystallization, so that the partial recrystallization region rolling is started in a state of uniform grains without coarse grains. If the reduction in the final rolling pass in the recrystallization temperature range is less than 10%, the recrystallization is insufficient, so that the grains grow into coarse grains during the holding time from rough rolling to the start of finish rolling. Therefore, the reduction in the final rolling pass in the recrystallization temperature range is 10% or more. The reduction in the final rolling pass in the recrystallization temperature range is preferably 11% or more. The reduction in the final rolling pass in the recrystallization temperature range is more preferably 13% or more, and even more preferably 15% or more. There is no particular limit to the upper limit of the reduction in the final rolling pass in the recrystallization temperature range, and the higher the reduction, the better. However, if it exceeds 70%, the productivity drops significantly, so it is preferable to make it 70% or less.
　［（再結晶温度－８０℃）以上における最終圧延パスの圧下率：１５％以上］
　再結晶域圧延完了後も部分的には再結晶するため、圧下率をさらに高めることで、再結晶の促進が可能であり、上位２０％粒径の微細化に有効である。よって、（再結晶温度－８０℃）以上における最終圧延パスの圧下率は１５％以上とする。（再結晶温度－８０℃）以上における最終圧延パスの圧下率は好ましくは１６％以上とする。（再結晶温度－８０℃）以上における最終圧延パスの圧下率は、より好ましくは１８％以上であり、さらに好ましくは２０％以上である。（再結晶温度－８０℃）以上における最終圧延パスの圧下率の上限は特に限定されなく、高いほど好ましいが、４０％を超えると生産性が著しく低下するため、４０％以下とすることが好ましい。 [Reduction rate of final rolling pass at (recrystallization temperature - 80 ° C) or higher: 15% or more]
Since partial recrystallization occurs even after the recrystallization region rolling is completed, further increasing the reduction ratio makes it possible to promote recrystallization and is effective in refining the top 20% grain size. Therefore, the reduction ratio of the final rolling pass at (recrystallization temperature - 80°C) or higher is set to 15% or more. The reduction ratio of the final rolling pass at (recrystallization temperature - 80°C) or higher is preferably set to 16% or more. The reduction ratio of the final rolling pass at (recrystallization temperature - 80°C) or higher is more preferably 18% or more, and even more preferably 20% or more. There is no particular upper limit to the reduction ratio of the final rolling pass at (recrystallization temperature - 80°C) or higher, and the higher the better, but since productivity drops significantly when the reduction ratio exceeds 40%, it is preferable to set the reduction ratio to 40% or less.
　（再結晶温度－８０℃）未満における圧延は、低温で圧延した方が、歪が多く導入されるため、結晶粒微細化に有効である。このため、制御冷却の冷却開始温度を遵守できる範囲内で、低温で圧延するのが好ましい。 Rolling at a temperature lower than (recrystallization temperature - 80°C) is more effective at refining the grains because more strain is introduced at lower temperatures. For this reason, it is preferable to roll at low temperatures within the range where the cooling start temperature for controlled cooling can be observed.
　熱間圧延工程において、結晶粒を小さくするためには、圧延終了温度は低いほどよいが、高圧水素環境下において耐ＨＩＳＣ性を確保する観点から、制御冷却の冷却開始温度が、熱延鋼板表面温度でＡｒ３点以上を確保できるよう圧延終了温度を設定する必要がある。ここで、Ａｒ３点とは、冷却中におけるフェライト変態開始温度を意味し、例えば、鋼の成分から以下の式で求めることができる。なお、熱延鋼板の表面温度は放射温度計等で測定することができる。
Ａｒ３（℃）＝９１０－３１０[％Ｃ]－８０[％Ｍｎ]－２０[％Ｃｕ]－１５[％Ｃｒ]－５５[％Ｎｉ]－８０[％Ｍｏ]
ただし、[％Ｘ]はＸ元素の鋼中含有量（質量％）を示す。 In the hot rolling process, in order to make the crystal grains smaller, the lower the rolling end temperature, the better. However, from the viewpoint of ensuring HISC resistance in a high-pressure hydrogen environment, it is necessary to set the rolling end temperature so that the cooling start temperature of the controlled cooling can ensure that the surface temperature of the hot-rolled steel sheet is equal to or higher than the Ar 3 point. Here, the Ar 3 point means the ferrite transformation start temperature during cooling, and can be calculated, for example, from the composition of the steel by the following formula. The surface temperature of the hot-rolled steel sheet can be measured by a radiation thermometer or the like.
Ar3 (℃) = 910-310[%C]-80[%Mn]-20[%Cu]-15[%Cr]-55[%Ni]-80[%Mo]
Here, [%X] indicates the content (mass%) of the X element in the steel.
　圧延後の冷却工程（制御冷却工程）
　［制御冷却の冷却開始温度：熱延鋼板の表面温度でＡｒ３変態点以上］
　冷却開始時の鋼板表面温度がＡｒ３変態点（Ａｒ３点）未満の場合、制御冷却前にフェライトが生成して、強度低下が大きくなる。このため、冷却開始時の熱延鋼板の表面温度はＡｒ３変態点以上とする。冷却開始時の熱延鋼板の表面温度はＡｒ３変態点＋２０℃以上とすることが好ましく、Ａｒ３変態点＋５０℃以上とすることがより好ましい。なお、冷却開始時の熱延鋼板の表面温度は、冷却開始温度が最も低くなる熱延鋼板尾端部の温度である。冷却開始時の熱延鋼板の表面温度はＡｒ３変態点＋１２０℃以下とすることが好ましく、Ａｒ３変態点＋８０℃以下とすることがより好ましい。 Cooling process after rolling (controlled cooling process)
[Controlled cooling start temperature: Surface temperature of hot-rolled steel sheet above Ar3 transformation point]
If the surface temperature of the steel sheet at the start of cooling is lower than the Ar3 transformation point ( Ar3 point), ferrite is generated before controlled cooling, and the strength is greatly reduced. Therefore, the surface temperature of the hot-rolled steel sheet at the start of cooling is set to be equal to or higher than the Ar3 transformation point. The surface temperature of the hot-rolled steel sheet at the start of cooling is preferably equal to or higher than the Ar3 transformation point + 20°C, and more preferably equal to or higher than the Ar3 transformation point + 50°C. The surface temperature of the hot-rolled steel sheet at the start of cooling is the temperature of the tail end of the hot-rolled steel sheet, where the cooling start temperature is the lowest. The surface temperature of the hot-rolled steel sheet at the start of cooling is preferably equal to or lower than the Ar3 transformation point + 120°C, and more preferably equal to or lower than the Ar3 transformation point + 80°C.
　［制御冷却の熱延鋼板先端と熱延鋼板尾端の冷却開始時間差：５０秒以内］
　冷却開始時の熱延鋼板圧延方向の先端と尾端の時間差が５０秒（ｓ）超えの場合、冷却開始時の先端と尾端の温度差が大きくなるため、冷却停止時の温度ばらつきが大きくなり、鋼材表面（鋼管の場合は鋼管内面の表面）から０．２５ｍｍにおけるビッカース硬さのばらつきが大きくなると共に耐ＨＩＳＣ性が劣化する。このため、熱延鋼板先端と熱延鋼板尾端の冷却開始時間差は５０秒以内とする。冷却開始時間差は好ましくは４５秒以内とする。冷却開始時間差は、より好ましくは４０秒以内とし、さらに好ましくは３２秒以内とする。熱延鋼板長が短くなることで冷却開始時間差を短くすることが可能であるが、製造性が低下するため、熱延鋼板搬送速度を速くすることで冷却開始時間差を短くすることが好ましい。冷却開始時間差は０秒であってよいが、製造性の観点から、２０秒以上とすることが好ましい。 [Time difference between the start of controlled cooling of the hot-rolled steel plate tip and tail: within 50 seconds]
When the time difference between the leading end and the tail end in the rolling direction of the hot-rolled steel sheet at the start of cooling exceeds 50 seconds (s), the temperature difference between the leading end and the tail end at the start of cooling becomes large, so that the temperature variation at the time of stopping cooling becomes large, and the variation in Vickers hardness at 0.25 mm from the steel surface (the surface of the inner surface of the steel pipe in the case of a steel pipe) becomes large and the HISC resistance deteriorates. For this reason, the cooling start time difference between the leading end and the tail end of the hot-rolled steel sheet is set to 50 seconds or less. The cooling start time difference is preferably set to 45 seconds or less. The cooling start time difference is more preferably set to 40 seconds or less, and further preferably set to 32 seconds or less. Although it is possible to shorten the cooling start time difference by shortening the length of the hot-rolled steel sheet, this reduces manufacturability, so it is preferable to shorten the cooling start time difference by increasing the hot-rolled steel sheet conveying speed. The cooling start time difference may be 0 seconds, but from the viewpoint of manufacturability, it is preferable to set it to 20 seconds or more.
　［制御冷却の平均冷却速度］
　板厚中央における７５０℃から５５０℃までの平均冷却速度：１５～５０℃／ｓ
　板厚中央における７５０℃から５５０℃までの平均冷却速度が１５℃／ｓ未満では、グラニュラーベイナイトを含む所定のベイナイト組織が得られずに強度低下が生じる。このため、板厚中央での平均冷却速度は１５℃／ｓ以上とする。組織のばらつき抑制の観点からは、板厚中央の平均冷却速度は１７℃／ｓ以上とすることが好ましい。板厚中央の平均冷却速度は２０℃／ｓ以上とすることが好ましく、２５℃／ｓ以上とすることがより好ましい。一方、ベイナイト組織の粒径のばらつきを抑制するために、平均冷却速度は５０℃／ｓ以下とする。平均冷却速度は４８℃／ｓ以下とすることが好ましく、４５℃／ｓ以下とすることがより好ましい。平均冷却速度は４２℃／ｓ以下とすることがさらに好ましく、３８℃／ｓ以下とすることがもっとも好ましい。なお、板厚中央における熱延鋼板温度で５５０℃以下の冷却については、特に限定されないが、組織や粒径のばらつき抑制の観点から、平均冷却速度は１５℃／ｓ以上５０℃／ｓ以下とすることが好ましい。つまり、５５０℃以下の冷却については、平均冷却速度は１５℃／ｓ以上とすることが好ましい。平均冷却速度は３０℃／ｓ以上とすることがより好ましく、３５℃／ｓ以上とすることがさらに好ましい。５５０℃以下の冷却については、平均冷却速度は５０℃／ｓ以下とすることが好ましい。平均冷却速度は４８℃／ｓ以下とすることがより好ましく、４２℃／ｓ以下とすることがさらに好ましい。５５０℃以下の平均冷却速度は５５０℃から２５０℃までの冷却速度の平均値である。 [Average cooling rate of controlled cooling]
Average cooling rate from 750 ° C to 550 ° C at the center of plate thickness: 15 to 50 ° C / s
If the average cooling rate from 750°C to 550°C at the center of the plate thickness is less than 15°C/s, the specified bainite structure including granular bainite cannot be obtained, and strength is reduced. For this reason, the average cooling rate at the center of the plate thickness is set to 15°C/s or more. From the viewpoint of suppressing the variation in the structure, the average cooling rate at the center of the plate thickness is preferably set to 17°C/s or more. The average cooling rate at the center of the plate thickness is preferably set to 20°C/s or more, and more preferably set to 25°C/s or more. On the other hand, in order to suppress the variation in the grain size of the bainite structure, the average cooling rate is set to 50°C/s or less. The average cooling rate is preferably set to 48°C/s or less, and more preferably set to 45°C/s or less. The average cooling rate is more preferably set to 42°C/s or less, and most preferably set to 38°C/s or less. The cooling to 550°C or less at the hot-rolled steel plate temperature at the center of the plate thickness is not particularly limited, but from the viewpoint of suppressing the variation in the structure and grain size, the average cooling rate is preferably set to 15°C/s or more and 50°C/s or less. That is, for cooling to 550°C or less, the average cooling rate is preferably 15°C/s or more. The average cooling rate is more preferably 30°C/s or more, and even more preferably 35°C/s or more. For cooling to 550°C or less, the average cooling rate is preferably 50°C/s or less. The average cooling rate is more preferably 48°C/s or less, and even more preferably 42°C/s or less. The average cooling rate for 550°C or less is the average value of the cooling rates from 550°C to 250°C.
　［冷却停止温度：２５０～６５０℃］
　熱間圧延後の板厚中央の冷却停止温度が６５０℃超えでは材料強度が大きく低下し、かつ均一なベイナイト組織を得る観点からも、板厚中央の冷却停止温度を６５０℃以下とする。板厚中央の冷却停止温度は６２０℃以下とすることが好ましく、６１５℃以下とすることがより好ましく、６００℃以下とすることがさらに好ましい。一方、板厚中央の冷却停止温度が２５０℃未満では、冷却時の焼割れが発生しやすくなる。また、均一なベイナイト組織を得るため、冷却停止温度は２５０℃以上とする。板厚中央の冷却停止温度は３００℃以上とすることが好ましく、３５０℃以上とすることがより好ましく、３８０℃以上とすることがさらに好ましい。鋼中水素量を抑制するという点からも冷却停止温度は所定の温度以上とする必要がある。具体的に、冷却中に鋼中に存在した水素は徐々に抜けていき、高温程その効果は大きいが、冷却停止温度が低すぎる場合には過冷却となり、鋼中に水素が残存する。さらに、冷却停止温度を低くしすぎると、他の相と比較して多量に水素を急増する残留オーステナイトが形成されやすくなる。そのため、鋼中水素量を低減させるためにも、冷却停止温度は２５０℃以上とする必要がある。冷却停止後は放冷すればよいが、ベイナイトの生成を促進するために、冷却停止温度から５０℃程度温度が下がるまでは徐冷することがより好ましい。 [Cooling stop temperature: 250 to 650°C]
If the cooling stop temperature at the center of the thickness after hot rolling exceeds 650°C, the material strength is significantly reduced, and from the viewpoint of obtaining a uniform bainite structure, the cooling stop temperature at the center of the thickness is set to 650°C or less. The cooling stop temperature at the center of the thickness is preferably set to 620°C or less, more preferably set to 615°C or less, and even more preferably set to 600°C or less. On the other hand, if the cooling stop temperature at the center of the thickness is less than 250°C, quench cracks are likely to occur during cooling. In addition, in order to obtain a uniform bainite structure, the cooling stop temperature is set to 250°C or more. The cooling stop temperature at the center of the thickness is preferably set to 300°C or more, more preferably set to 350°C or more, and even more preferably set to 380°C or more. In order to suppress the amount of hydrogen in the steel, the cooling stop temperature needs to be set to a predetermined temperature or more. Specifically, hydrogen present in the steel gradually escapes during cooling, and the higher the temperature, the greater the effect, but if the cooling stop temperature is too low, the steel will be overcooled and hydrogen will remain in the steel. Furthermore, if the cooling stop temperature is too low, it is easy to form retained austenite, which rapidly increases hydrogen in a large amount compared to other phases. Therefore, in order to reduce the amount of hydrogen in the steel, the cooling stop temperature needs to be 250°C or higher. Although it is acceptable to allow the steel to cool after cooling is stopped, it is more preferable to cool it slowly until the temperature drops by about 50°C from the cooling stop temperature in order to promote the formation of bainite.
　［脱水素処理（好適条件）］
　鋼材中にそもそも水素が存在する場合には疲労き裂進展の加速が増大され、疲労寿命が低下する。そのため、製造後に残存する水素を放出させるために、脱水素処理を用いることが好ましい。脱水素処理は、製品使用前に高温で一定時間保持することで鋼中水素量を低減させることができる。
また、水素は室温においても長時間保持することで脱水素処理が可能である。室温で保持する際は保持時間が長期間化するため、保持時間は９６時間以上が好ましい。さらに、鋼表面のスケールは脱水素を阻害するため、スケールを除去し脱水素処理行う方が好ましい。保持時間Ｒ（ｓｅｃ）は、鋼板および鋼管の板厚並びに管厚ｔ（ｍｍ）、および室温における鋼中の水素拡散係数Ｄ（ｍｍ・ｓｅｃ－１）から、以下の式（Ａ）とすることが好ましい。
Ｒ≧ｔ２／Ｄ・・・（Ａ）
水素拡散係数は含有している成分や金属組織によっても変わるが、例えば、水素拡散係数は１×１０－５～５×１０－３ｍｍ２／ｓを採用しても良い。より好ましくは５×１０－４ｍｍ２／ｓ以下である。 
脱水素処理工程は、造管または鋼管をつなげる溶接施工前に実施する。なお、脱水素処理は高温の水素拡散係数Ｄが小さくなり、早く水素が抜けるため高温である方が好ましい。
高温の場合は上記（Ａ）式のＤの値を保持する温度の拡散係数Ｄ’（それぞれの温度における拡散係数）を用いて計算しても良い。一方、脱水素工程の温度Ｔが高すぎる場合には材料強度が著しく低下するため、脱水素処理温度は５５０℃以下が好ましい。脱水素処理温度Ｔは５００℃以下とすることがより好ましい。脱水素処理温度Ｔは４００℃以下とすることがさらに好ましく、３００℃以下とすることがもっとも好ましい。また、室温よりも温度を低下させた脱水素処理は処理時間およびコスト増の要因であるという理由から脱水素処理温度Ｔは室温以上とすることが好ましい。脱水素処理温度Ｔは５０℃以上とすることがより好ましい。脱水素処理温度Ｔは１００℃以上とすることがさらに好ましく、１５０℃以上とすることがもっとも好ましい。ここで述べている脱水素処理温度Ｔとは脱水素処理工程における雰囲気の温度である。室温とは２０±１０℃のことをいう。 [Dehydrogenation treatment (optimal conditions)]
If hydrogen is present in steel to begin with, the acceleration of fatigue crack growth increases, and the fatigue life decreases. Therefore, it is preferable to use a dehydrogenation process to release the hydrogen remaining after manufacturing. Dehydrogenation can reduce the amount of hydrogen in steel by holding the product at high temperature for a certain period of time before use.
Hydrogen can also be dehydrogenated by holding it at room temperature for a long time. When holding it at room temperature, the holding time is long, so the holding time is preferably 96 hours or more. Furthermore, since scale on the steel surface inhibits dehydrogenation, it is preferable to remove the scale before dehydrogenation. The holding time R (sec) is preferably calculated from the plate thickness and pipe thickness t (mm) of the steel plate and steel pipe, and the hydrogen diffusion coefficient D (mm·sec −1 ) in steel at room temperature, as shown in the following formula (A).
R≧ t2 /D (A)
The hydrogen diffusion coefficient varies depending on the contained components and metal structure, but for example, the hydrogen diffusion coefficient may be 1×10 −5 to 5×10 −3 mm 2 /s, and more preferably 5×10 −4 mm 2 /s or less.
The dehydrogenation process is carried out before pipe making or welding to connect steel pipes. It is preferable to carry out the dehydrogenation process at a high temperature because the hydrogen diffusion coefficient D at high temperatures becomes small and hydrogen is quickly removed.
In the case of high temperatures, the diffusion coefficient D' (diffusion coefficient at each temperature) at which the value of D in the above formula (A) is maintained may be used for calculation. On the other hand, if the temperature T of the dehydrogenation process is too high, the material strength is significantly reduced, so the dehydrogenation temperature is preferably 550°C or less. It is more preferable that the dehydrogenation temperature T is 500°C or less. It is even more preferable that the dehydrogenation temperature T is 400°C or less, and most preferably 300°C or less. In addition, since dehydrogenation at a temperature lower than room temperature is a factor that increases the processing time and cost, it is preferable that the dehydrogenation temperature T is room temperature or higher. It is more preferable that the dehydrogenation temperature T is 50°C or higher. It is more preferable that the dehydrogenation temperature T is 100°C or higher, and most preferably 150°C or higher. The dehydrogenation temperature T mentioned here is the temperature of the atmosphere in the dehydrogenation process. Room temperature means 20±10°C.
　特に、加熱する場合、鋼材および鋼管の板厚中央の温度Ｔｃが脱水素処理工程における雰囲気の温度（脱水素処理温度Ｔ）に到達するまでに時間を要するため、雰囲気温度において上記保持時間Ｒ（ｓｅｃ）を満たしていても、板厚中央が脱水素処理温度Ｔ（雰囲気温度）に達していない場合は脱水素処理が不十分となる可能性がある。そのため、板厚中央温度Ｔｃが目標とする脱水素処理温度Ｔに達してからＲ時間（ｓｅｃ）以上保持することが好ましい。さらに、所定の水素ガス中の水素中破壊靭性を得るために、表層部と板厚中央の鋼材水素量を適切に調整する必要があり、そのために、脱水素処理温度Ｔ（雰囲気温度）で、（Ａ）式で規定されたＲ（ｓｅｃ）以上保持することが好ましく、さらに板厚中央温度Ｔｃが目標とする脱水素処理温度Ｔに達してから上記保持時間Ｒ（ｓｅｃ）以上保持することがより好ましい。言い換えると、少なくとも前者は鋼材および鋼管の表層部の鋼材水素量を適切に制御でき、後者まで実施すると鋼材および鋼管の表層部から板厚中央までの鋼材水素量を適切に制御することができる。板厚温度が板厚中央温度Ｔｃは熱電対などを用いて実測してもいいし、有限要素法などを用いて予測してもよい。 In particular, when heating, it takes time for the temperature Tc at the center of the thickness of the steel material and steel pipe to reach the temperature of the atmosphere in the dehydrogenation process (dehydrogenation temperature T), so even if the above-mentioned holding time R (sec) is met at the atmospheric temperature, if the center of the thickness does not reach the dehydrogenation temperature T (atmospheric temperature), the dehydrogenation may be insufficient. Therefore, it is preferable to hold the temperature Tc for R time (sec) or more after it reaches the target dehydrogenation temperature T. Furthermore, in order to obtain hydrogen fracture toughness in a specified hydrogen gas, it is necessary to appropriately adjust the hydrogen content of the steel material in the surface layer and the center of the thickness. For this reason, it is preferable to hold the temperature Tc at the dehydrogenation temperature T (atmospheric temperature) for R (sec) or more as specified by formula (A), and it is even more preferable to hold the temperature Tc for the above-mentioned holding time R (sec) or more after it reaches the target dehydrogenation temperature T. In other words, at least the former can appropriately control the amount of hydrogen in the surface layer of the steel material and steel pipe, and by implementing the latter, the amount of hydrogen in the steel material from the surface layer to the center of the thickness of the steel material and steel pipe can be appropriately controlled. The thickness temperature, or the center temperature Tc, can be measured using a thermocouple or the like, or can be predicted using the finite element method or the like.
　なお、脱水素処理工程の時間と温度は、後述しているとおり電縫管やＵＯＥ等の造管工程で加熱する際に加えられた温度と時間が含まれても良い。さらに、鋼表面のスケールは脱水素を阻害するため、スケールを除去し脱水素処理を行う方が好ましい。除去方法は問わないが、例えば高圧洗浄による物理的な洗浄でもよいし、スケール除去剤を用いた化学的な手法を用いてもよい。厚みとして１００μｍ程度除去されればスケール除去の効果が得られる。 The time and temperature of the dehydrogenation process may include the temperature and time applied when heating in the pipe-making process for electric resistance welded pipes, UOE, etc., as described below. Furthermore, since scale on the steel surface inhibits dehydrogenation, it is preferable to remove the scale before carrying out the dehydrogenation process. There is no restriction on the removal method, but it may be physical cleaning using a high-pressure cleaner, for example, or a chemical method using a scale remover. The effect of scale removal can be obtained if a thickness of about 100 μm is removed.
　第２実施形態
　さらに、高強度ラインパイプ用鋼管の一例として挙げられるＵＯＥ鋼管は下記に示す製造条件を限定することにより得ることができ、製造方法および条件を具体的に説明する。ＵＯＥ鋼管の成分組成、金属組織、水素誘起き裂進展下限界ＫＩＨは第１実施形態の鋼板で説明した内容と同様であり、製造方法についても溶鋼工程、加熱工程、熱間圧延工程、熱間圧延後の制御冷却工程、脱水素処理工程は鋼材で説明した内容と同等の内容で実施される。下記では、圧延後の造管工程を具体的に説明する。 Second embodiment Furthermore, a UOE steel pipe, which is an example of a steel pipe for high strength line pipe, can be obtained by limiting the manufacturing conditions shown below, and the manufacturing method and conditions will be specifically described. The chemical composition, metal structure, and hydrogen-induced crack propagation lower limit K IH of the UOE steel pipe are the same as those described for the steel plate of the first embodiment, and as for the manufacturing method, the molten steel process, heating process, hot rolling process, controlled cooling process after hot rolling, and dehydrogenation process are performed in the same manner as those described for the steel material. The pipe-making process after rolling will be specifically described below.
　造管工程
　ＵＯＥ鋼管は、熱延鋼板を曲げ加工、具体的にいうと熱延鋼板の端部を開先加工し、Ｃプレス、Ｕプレス、Ｏプレスで鋼管形状に成形する加工を施した後、内面溶接および外面溶接で突き合わせ部をシーム溶接し、さらに必要に応じて拡管工程を経て製造される。また、溶接方法は十分な継手強度と継手靭性が得られる方法であれば、いずれの方法でも良いが、優れた溶接品質と製造能率の観点から、サブマージアーク溶接を用いることが好ましい。また、プレスベンド成形により管状に成形した後、突き合せ部をシーム溶接した鋼管に対しても、拡管を実施することができる。さらに、造管後の溶接熱影響部で上記の介在物が存在する場合、母材部と同様に水素集積源として作用し、耐ＨＩＣ性およびＫＩＨの劣化を引き起こす。溶接部介在物の低減にもＳあるいはＯの含有量を低減することも有効である。このため、溶接後の鋼管の１５００℃から１０００℃までの温度域における平均冷却速度を５０℃／分以上とすることが好ましい。平均冷却速度は５５℃／分以上とすることがより好ましく、６０℃／分以上とすることがさらに好ましい。上限は特に限定されるものではないが、平均冷却速度は１００℃／分以下とすることが好ましい。 Pipe-making process UOE steel pipes are manufactured by bending hot-rolled steel sheets, specifically by groove-forming the ends of the hot-rolled steel sheets, and forming them into a steel pipe shape using a C press, a U press, and an O press, then seam-welding the butt joints using internal and external welding, and then expanding the pipe as necessary. Any welding method may be used as long as it provides sufficient joint strength and joint toughness, but it is preferable to use submerged arc welding from the viewpoint of excellent welding quality and manufacturing efficiency. Pipe expansion can also be performed on steel pipes that have been formed into a tubular shape by press bending and then have seam-welded butt joints. Furthermore, if the above-mentioned inclusions exist in the welded heat-affected zone after pipe-making, they act as a hydrogen accumulation source in the same way as in the base metal, causing deterioration of HIC resistance and KIH. Reducing the content of S or O is also effective in reducing inclusions in the welded zone. For this reason, it is preferable to set the average cooling rate of the welded steel pipe in the temperature range from 1500°C to 1000°C to 50°C/min or more. The average cooling rate is more preferably 55° C./min or more, and even more preferably 60° C./min or more. Although there is no particular upper limit, the average cooling rate is preferably 100° C./min or less.
　第３実施形態
　さらに、本発明に係る高強度ラインパイプ用鋼管には、一例として電縫鋼管が挙げられ、電縫鋼管は下記に示す製造条件を限定することにより得ることができ、製造方法および条件を具体的に説明する。鋼材の成分組成、金属組織、水素誘起き裂進展下限界ＫＩＨは第１実施形態の鋼材で説明した内容と同様であり、製造方法についても圧延後の冷却工程、造管工程以外の工程（溶鋼工程、加熱工程、熱間圧延工程、脱水素処理工程）は鋼材で説明した内容と同等の内容で実施される。 Third embodiment Furthermore, an example of a steel pipe for high strength line pipe according to the present invention is an electric resistance welded steel pipe, which can be obtained by limiting the manufacturing conditions shown below, and the manufacturing method and conditions will be specifically described below. The composition, metal structure, and hydrogen induced crack propagation lower limit K IH of the steel material are the same as those described for the steel material of the first embodiment, and as for the manufacturing method, the steps other than the cooling step after rolling and the pipe making step (melting step, heating step, hot rolling step, dehydrogenation treatment step) are performed in the same manner as those described for the steel material.
　圧延後の冷却工程（制御冷却工程）
　制御冷却の冷却開始温度、制御冷却の平均冷却速度は第１実施形態で記載と同じ内容で実施される。 Cooling process after rolling (controlled cooling process)
The cooling start temperature and the average cooling rate of the controlled cooling are the same as those described in the first embodiment.
　［冷却停止温度：２５０～６５０℃］
　熱間圧延後の板厚中央の冷却停止温度が６５０℃超えでは材料強度が大きく低下し、かつ均一なベイナイト組織を得る観点からも、板厚中央の冷却停止温度を６５０℃以下とする。板厚中央の冷却停止温度は６２０℃以下とすることが好ましく、６１５℃以下とすることがより好ましく、６００℃以下とすることがさらに好ましい。一方、板厚中央の冷却停止温度が２５０℃未満では、冷却時の焼割れが発生しやすくなる。このため、板厚中央の冷却停止温度は２５０℃以上とする。板厚中央の冷却停止温度は３００℃以上とすることが好ましく、３５０℃以上とすることがより好ましく、３８０℃以上とすることがさらに好ましい。鋼板表面における硬質な組織の生成を確実に抑制するため、板厚中央の冷却停止温度は４５０℃以上がもっとも好ましい。冷却停止後は放冷すればよいが、ベイナイトの生成を促進するために、冷却停止温度から５０℃程度温度が下がるまでは徐冷することがより好ましい。 [Cooling stop temperature: 250 to 650°C]
If the cooling stop temperature at the center of the thickness after hot rolling exceeds 650°C, the material strength is significantly reduced, and from the viewpoint of obtaining a uniform bainite structure, the cooling stop temperature at the center of the thickness is set to 650°C or less. The cooling stop temperature at the center of the thickness is preferably set to 620°C or less, more preferably set to 615°C or less, and even more preferably set to 600°C or less. On the other hand, if the cooling stop temperature at the center of the thickness is less than 250°C, quench cracks are likely to occur during cooling. For this reason, the cooling stop temperature at the center of the thickness is set to 250°C or more. The cooling stop temperature at the center of the thickness is preferably set to 300°C or more, more preferably set to 350°C or more, and even more preferably set to 380°C or more. In order to reliably suppress the generation of hard structures on the steel sheet surface, the cooling stop temperature at the center of the thickness is most preferably 450°C or more. After cooling is stopped, it is sufficient to allow the steel to cool, but in order to promote the generation of bainite, it is more preferable to slowly cool the steel until the temperature drops by about 50°C from the cooling stop temperature.
　その後、上記のようにして得られた熱延鋼板をコイル状に巻取る。巻取り温度は６５０℃以下とすることが好ましい。巻取り温度は６２０℃以下とすることがより好ましく、６１５℃以下とすることがより好ましく、６００℃以下とすることがさらに好ましい。下限は、巻取り温度は２５０℃以上とすることが好ましく、３００℃以上とすることがより好ましい。３５０℃以上とすることがさらに好ましく、３８０℃以上とすることがもっとも好ましい。 Then, the hot-rolled steel sheet obtained as described above is wound into a coil. The winding temperature is preferably 650°C or less. The winding temperature is more preferably 620°C or less, more preferably 615°C or less, and even more preferably 600°C or less. As for the lower limit, the winding temperature is preferably 250°C or more, more preferably 300°C or more, more preferably 350°C or more, and most preferably 380°C or more.
　造管工程
　本発明の一例として挙げている電縫鋼管は、冷間ロール成形により円筒状に成形し、該円筒状の周方向両端部を突き合わせて溶接することによって製造される。さらに、以下の（１）式を満たすサイジングロールを用いて電縫鋼管素材（電縫鋼管）に成形し（サイジング工程）、前記電縫鋼管素材の内面に以下の（２）式を満たす内圧ｐ（ＭＰａ）を負荷する（内圧負荷工程）ことによって製造してもよい。
　なお、前記円筒状とは、管周断面が「Ｃ」形状であることを指す。 Pipe-making process The electric resistance welded steel pipe given as an example of the present invention is manufactured by forming the pipe into a cylindrical shape by cold roll forming, and then butting and welding both circumferential ends of the cylindrical shape together. Furthermore, the electric resistance welded steel pipe may be manufactured by forming the pipe into an electric resistance welded steel pipe material (electric resistance welded steel pipe) using a sizing roll that satisfies the following formula (1) (sizing process), and applying an internal pressure p (MPa) that satisfies the following formula (2) to the inner surface of the electric resistance welded steel pipe material (internal pressure application process).
The cylindrical shape means that the circumferential cross section of the tube is in a "C" shape.
　サイジングロールの直径（ｍｍ）≧熱延鋼板の板厚（ｍｍ）／０．０２０　・・・（１）
熱延鋼板の板厚とは、サイジング工程を行う前の熱延鋼板の板厚のことである。
　Ｘ＜ｐ≦Ｘ×１．５　・・・（２）
なお、Ｘ＝（電縫鋼管素材の肉厚（ｍｍ）／電縫鋼管素材の半径（ｍｍ））×電縫鋼管素材の降伏強度（ＭＰａ）
前記した内圧の負荷は、例えば、ゴム素材のパッキンで管端を封じて管内部に水圧を負荷することにより実施することができる。また、形状を安定化させるために、必要に応じて外枠として所期した径の金型を使用することもできる。 Diameter of sizing roll (mm) ≧ Thickness of hot-rolled steel sheet (mm)/0.020 (1)
The plate thickness of the hot-rolled steel plate means the plate thickness of the hot-rolled steel plate before the sizing process is performed.
X<p≦X×1.5 (2)
In addition, X = (thickness of electric welded steel pipe material (mm) / radius of electric welded steel pipe material (mm)) × yield strength of electric welded steel pipe material (MPa)
The above-mentioned internal pressure can be applied, for example, by sealing the end of the tube with a rubber packing and applying water pressure to the inside of the tube. In order to stabilize the shape, a mold of a desired diameter can be used as an outer frame as necessary.
　なお、本発明の鋼管の一例として挙げている電縫鋼管素材の肉厚は５ｍｍ以上が好ましい。電縫鋼管素材の肉厚は３０ｍｍ以下が好ましい。電縫鋼管素材の半径は、上限は規定しないが、大きくなると設備の負荷が増大するため、電縫管素材の半径は、４００ｍｍ以下が好ましい。電縫管素材の半径は、２００ｍｍ以上が好ましい。また、電縫鋼管素材の降伏強度は、パイプライン操業ガス圧力に耐えるため、４８０ＭＰａ以上が好ましい。降伏強度は５００ＭＰａ以上がより好ましい。一方、水素脆化感受性増大を避けるために、降伏強度は５６０ＭＰａ以下が好ましい。降伏強度は５５０ＭＰａ以下がより好ましい。 The thickness of the electric welded steel pipe material given as an example of the steel pipe of the present invention is preferably 5 mm or more. The thickness of the electric welded steel pipe material is preferably 30 mm or less. There is no upper limit for the radius of the electric welded steel pipe material, but since a larger radius increases the load on the equipment, the radius of the electric welded pipe material is preferably 400 mm or less. The radius of the electric welded pipe material is preferably 200 mm or more. Furthermore, the yield strength of the electric welded steel pipe material is preferably 480 MPa or more in order to withstand the gas pressure of pipeline operation. A yield strength of 500 MPa or more is more preferable. On the other hand, in order to avoid increased susceptibility to hydrogen embrittlement, the yield strength is preferably 560 MPa or less. A yield strength of 550 MPa or less is more preferable.
　サイジング工程では、ロール通過時にロール形状に沿って管軸方向に曲げ変形が生じ、管軸方向の残留応力が発生する。前記曲げ変形における曲げひずみが大きいほど、管軸方向の残留応力の絶対値が大きくなる。前記曲げひずみは、サイジングロールの直径が小さいほど、また熱延鋼板の板厚が大きいほど大きくなる。
よって、本発明では、せん断残留応力を低くする観点から、管軸方向の残留応力の絶対値を小さくするため、サイジングロールの直径は前記（１）式を満足させるものとする。
サイジングロールの直径が前記（１）式の右辺未満の場合、本発明で目的とするせん断残留応力が得られない。なお、特にサイジングロールの直径の上限は規定しないが、サイジングロールが大きくなると設備の負荷が増大するため、サイジングロールの直径は２０００ｍｍ以下とすることが好ましい。 In the sizing process, bending deformation occurs in the tube axial direction along the roll shape when the steel sheet passes through the rolls, and residual stress in the tube axial direction is generated. The larger the bending strain in the bending deformation, the larger the absolute value of the residual stress in the tube axial direction. The bending strain increases as the diameter of the sizing rolls decreases and as the thickness of the hot-rolled steel sheet increases.
Therefore, in the present invention, from the viewpoint of reducing the residual shear 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 axial direction of the tube.
If the diameter of the sizing roll is less than the right side of the formula (1), the intended residual shear stress of the present invention cannot be obtained. Although there is no particular upper limit to the diameter of the sizing roll, the larger the sizing roll, the greater the load on the equipment, so that the diameter of the sizing roll is preferably 2000 mm or less.
　内圧負荷工程では、電縫鋼管素材を拡管することにより、管周方向に引張応力を発生させて、管周方向の残留応力の絶対値を小さくする。
かかる内圧負荷工程の内圧ｐ（ＭＰａ）が大きいほど、管周方向の残留応力の絶対値が小さくなる。管周方向に発生する引張応力は、鋼管の半径が大きいほど、鋼管の肉厚が小さいほど、高くなる。 In the internal pressure application process, the electric resistance welded steel pipe material is expanded to generate tensile stress in the circumferential direction of the pipe, thereby reducing the absolute value of the residual stress in the circumferential direction of the pipe.
The greater the internal pressure p (MPa) in the internal pressure loading step, the smaller the absolute value of the residual stress in the circumferential direction of the pipe. 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 circumferential direction of the pipe.
　前記（２）式の左辺（Ｘ）は、管周方向に発生する引張応力が電縫鋼管素材の降伏応力に等しくなる場合の内圧ｐに対応する。
本発明では、せん断残留応力を低くする観点から、管軸方向の残留応力の絶対値を小さくするため、内圧ｐを（２）式の左辺（Ｘ）より大きい値とし、電縫鋼管素材を塑性域まで拡管させる。一方、内圧ｐが（２）式の右辺（Ｘ×１．５）超になると、管周方向の残留応力の絶対値は小さくなるが、拡管による加工硬化量が大きくなり過ぎて、管表面の転位密度が上昇し、水素中の破壊靭性が低下する。 The left side (X) of the above equation (2) corresponds to the internal pressure p when the tensile stress generated in the circumferential direction of the pipe is equal to the yield stress of the electric resistance welded steel pipe material.
In the present invention, in order to reduce the absolute value of the residual stress in the pipe axial direction from the viewpoint of reducing the shear residual stress, the internal pressure p is set to a value greater than the left side (X) of equation (2) and the electric resistance welded steel pipe material is expanded to the plastic region. On the other hand, when the internal pressure p exceeds the right side (X × 1.5) of equation (2), the absolute value of the residual stress in the pipe circumferential direction decreases, but the amount of work hardening due to pipe expansion becomes too large, the dislocation density on the pipe surface increases, and the fracture toughness in hydrogen decreases.
　一部は上記で説明しているとおり、高強度鋼管については、本開示の高強度鋼材を、プレスベンド成形、ロール成形、ＵＯＥ成形等で管状に成形した後、突き合わせ部を溶接することにより、原油や天然ガスの輸送に好適な鋼板内の材質均一性に優れた耐サワーラインパイプ用高強度鋼管（ＵＯＥ鋼管、電縫鋼管、スパイラル鋼管等）を製造することができる。また、本開示の高強度鋼板を鋼管に用いることにより、溶接部の高硬度域が存在しても、耐ＨＩＳＣ性に優れる鋼管を製造することができる。 As partially explained above, for high-strength steel pipes, the high-strength steel material disclosed herein can be formed into a tube by press bending, roll forming, UOE forming, or the like, and then the butt joints can be welded to produce high-strength steel pipes for sour-resistant line pipes (UOE steel pipes, electric resistance welded steel pipes, spiral steel pipes, etc.) with excellent material uniformity within the steel plate, suitable for transporting crude oil or natural gas. In addition, by using the high-strength steel plate disclosed herein for steel pipes, steel pipes with excellent HISC resistance can be produced, even if there is a high hardness region in the weld.
　実施例に基づいて以下で本発明をさらに具体的に説明する。以下の実施例は、本発明の好適な一例を示すものであり、本発明は、以下の実施例によって何ら限定されるものではない。 The present invention will be described in more detail below with reference to examples. The following examples are intended to illustrate preferred embodiments of the present invention, and the present invention is not limited in any way by the following examples.
　表１－１、１－２に示す成分組成のスラブを作製し、前記スラブを熱間圧延し、制御冷却を行い、脱水素処理を実施し、鋼材を得た。またその鋼材を成形して鋼管を製造した。製造条件は表２－１、２－２に示す。
Ｎｏ．２～６、１２～２２、３５、３７は得られた鋼材（熱延鋼板）を曲げ加工し、両端部を突合せて溶接する造管工程を実施し、Ｎｏ．７～１１、２３～３３、３６、３８は得られた鋼材（熱延鋼板）を冷間ロール成形により円筒状に成形し、該円筒状の周方向両端部を突合せて電縫溶接する造管工程を実施して、鋼管成形した。Ｎｏ．１、３４は鋼材ままとした。
得られた鋼材と鋼管のそれぞれについて、金属組織および材質を評価した結果を表３－１、３－２に示す。評価方法は、以下の通りである。 Slabs having the composition shown in Tables 1-1 and 1-2 were prepared, and the slabs were hot-rolled, controlled cooled, and dehydrogenated to obtain steel materials. The steel materials were then formed to produce steel pipes. The manufacturing conditions are shown in Tables 2-1 and 2-2.
For Nos. 2 to 6, 12 to 22, 35, and 37, the obtained steel material (hot-rolled steel sheet) was bent and both ends were butted together and welded in a pipe-making process, while for Nos. 7 to 11, 23 to 33, 36, and 38, the obtained steel material (hot-rolled steel sheet) was formed into a cylindrical shape by cold roll forming, and both circumferential ends of the cylindrical shape were butted together and electric resistance welded in a pipe-making process to form a steel pipe. For Nos. 1 and 34, the steel material was left as it was.
The metal structure and properties of the obtained steel material and steel pipe were evaluated, and the results are shown in Tables 3-1 and 3-2. The evaluation methods are as follows.
　残留オーステナイト測定
　上記に従って得られた鋼材および鋼管の長手方向中央部の板幅中央部より金属組織観察用サンプルを採取し、長手方向と平行な断面を観察対象面としてバフ研磨まで行い、その後、ピクリン酸エッチングにより表層を化学研磨により除去し、Ｘ線回折測定を用いて測定した。具体的に、入射Ｘ線にはＣｏ－Ｋα線源を用い、フェライトの（２００）、（２１１）、（２２０）面とオーステナイトの（２００）、（２２０）、（３１１）面の強度比から残留オーステナイトの面積分率を算出した。 Measurement of retained austenite Samples for metallographic observation were taken from the longitudinal center of the steel material and steel pipe obtained as described above, and the cross section parallel to the longitudinal direction was subjected to buffing and then the surface layer was removed by chemical polishing using picric acid etching, and the area fraction of retained austenite was calculated from the intensity ratio of the (200), (211), and (220) planes of ferrite to the (200), (220), and (311) planes of austenite.
　ベイナイトの最大粒径と面積分率の算出
　上記に従って得られた鋼材および鋼管の板幅中央部より金属組織観察用サンプルを採取し、このサンプルについて圧延長手方向と平行な断面を観察対象面とした。観察対象面を鏡面研磨したあと、コロイダルシリカでエッチングを行い、サンプル中央の位置で１ｍｍ×１ｍｍの視野でＥＢＳＤ（Ｅｌｅｃｔｒｏｎ　Ｂａｃｋｓｃａｔｔｅｒ　Ｄｉｆｆｒａｃｔｉｏｎ）法にて結晶データを収集した（測定ステップ：０．８μｍ）。結晶粒径の定義はＡｒｅａ粒径（方位差が１５°以上の境界を粒界と定義したときの加重平均）とした。上述の結晶データから各結晶粒径を求め、最大粒径を求めた。
また、面積分率については上記の観察対象面について、３ｖｏｌ％ナイタール溶液を用いてエッチングし、１０００～５０００倍間の適切な倍率で走査電子顕微鏡（ｓｃａｎｎｉｎｇ　ｅｌｅｃｔｒｏｎ　ｍｉｃｒｏｓｃｏｐｅ）写真を撮影し、ベイナイトを観察した。ベイナイトは、非特許文献１の組織写真と比較して目視で判断し、組織分率は、上記判断を基にＳＥＭ写真においてベイナイトとその他の領域を二値化して、画像解析（ｉｍａｇｅ　ａｎａｌｙｓｉｓ）により分率を求め、ベイナイトの面積分率とした。 Calculation of maximum grain size and area fraction of bainite Samples for metal structure observation were taken from the center of the plate width of the steel material and steel pipe obtained according to the above, and the cross section of this sample parallel to the rolling direction was used as the observation surface. After mirror polishing the observation surface, etching was performed with colloidal silica, and crystal data was collected by EBSD (Electron Backscatter Diffraction) method in a field of view of 1 mm x 1 mm at the center position of the sample (measurement step: 0.8 μm). The crystal grain size was defined as the area grain size (weighted average when the boundary with an orientation difference of 15° or more is defined as the grain boundary). Each crystal grain size was obtained from the above crystal data, and the maximum grain size was obtained.
Regarding the area fraction, the above-mentioned observation surface was etched using a 3 vol% nital solution, and a scanning electron microscope photograph was taken at an appropriate magnification between 1000 and 5000 times to observe bainite. Bainite was judged visually by comparison with the structure photograph in Non-Patent Document 1, and the structure fraction was determined by binarizing the bainite and other regions in the SEM photograph based on the above judgment, and determining the fraction by image analysis, which was taken as the area fraction of bainite.
　介在物の観察、個数密度の算出
　上記に従って得られた鋼材および鋼管の長手方向中央部の板幅中央部より金属組織観察用サンプルを採取し、このサンプルについて圧延長手方向と平行な断面を観察対象面とした。観察対象面を鏡面研磨したあと、コロイダルシリカでエッチングを行い、サンプル中央の位置で１０ｍｍ×１０ｍｍの視野で走査型電子顕微鏡（ＳＥＭ）による観察を実施した。観察倍率は２０００～５０００倍として観察し、３視野の平均を介在物の個数密度とした。 Observation of inclusions and calculation of number density Samples for metal structure observation were taken from the center of the plate width in the longitudinal center of the steel material and steel pipe obtained as described above, and the cross section of this sample parallel to the rolling direction was used as the observation surface. The observation surface was mirror-polished and then etched with colloidal silica, and observation was performed with a scanning electron microscope (SEM) at the center of the sample in a field of view of 10 mm x 10 mm. The observation was performed at a magnification of 2000 to 5000 times, and the average of three fields of view was used as the number density of inclusions.
　引張強さ（ＴＳ）
　上記に従って得られた鋼材および鋼管から、ＪＩＳ　Ｚ　２２０１に準拠してＪＩＳ１４号比例試験片（平行部直径７ｍｍ、標点間距離３５ｍｍ）を採取し、引張強さを測定した。 Tensile strength (TS)
From the steel materials and steel pipes obtained as described above, JIS No. 14 proportional test pieces (parallel part diameter 7 mm, gauge length 35 mm) were taken in accordance with JIS Z 2201, and the tensile strength was measured.
　水素昇温分析
　鋼中に残存する水素量は昇温脱離分析法を用いて、低温型昇温式水素分析装置〈ガスクロマトグラフタイプ〉（ＪＴＦ－２０ＡＬ）を用いた。昇温脱離分析は２００℃／ｈの昇温速度で室温から４００℃までの温度範囲で行い、その総和を水素量とした。試験体は鋼材の板厚１／４位置および鋼管の内面から１／４位置で鋼管長手方向に３０ｍｍ長さで直径７Φの円柱形状である。なお、この水素量は後述している時効で説明する高圧水素疲労試験に供する前であり、表１－１、１－２に示すＨ量である。 Temperature-programmed hydrogen analysis The amount of hydrogen remaining in the steel was measured using a temperature-programmed desorption analysis method, using a low-temperature temperature-programmed hydrogen analyzer (gas chromatograph type) (JTF-20AL). Temperature-programmed desorption analysis was performed in the temperature range from room temperature to 400°C at a heating rate of 200°C/h, and the sum of the measurements was taken as the amount of hydrogen. The test specimens were cylindrical, 30 mm long in the longitudinal direction of the steel pipe, at a 1/4 position of the plate thickness of the steel material and a 1/4 position from the inner surface of the steel pipe, and had a diameter of 7Φ. This amount of hydrogen was measured before the steel was subjected to the high-pressure hydrogen fatigue test described in the aging section below, and is the amount of H shown in Tables 1-1 and 1-2.
　高圧水素ガス中破壊靭性試験
　室温（２０±１０℃）、圧力２５ＭＰａの水素ガス（１００％水素を含む）、または上記温度、圧力を有する水素水素分圧として１ＭＰａ以上の水素を含む天然ガス（主成分はメタン、エタンなどの炭化水素）混合雰囲気中で、ＡＳＴＭ　Ｅ１８２０に準拠して実施した。試験片にはＣＴ試験片（板厚１２．７ｍｍ、板幅２５．４ｍｍ）を用い、機械ノッチ導入方向と鋼材の圧延方向が平行となる向きで採取した。疲労予き裂は大気中で導入し、周波数：１Ｈｚ、繰り返し荷重波形：正弦波、制御方法：Ｋ値制御、応力比Ｒ：０．１の条件とした。その後、水素ガス、または水素ガス＋天然ガス混合雰囲気下とした。破壊靭性試験は、単一試験片による除荷―弾性コンプライアンス法で実施した。荷重負荷時のクロスヘッド変位速度は０．００２ｍｍ／ｓｅｃとした。 Fracture toughness test in high pressure hydrogen gas The test was carried out in accordance with ASTM E1820 at room temperature (20±10°C), in hydrogen gas (containing 100% hydrogen) at a pressure of 25 MPa, or in a mixed atmosphere of natural gas (mainly hydrocarbons such as methane and ethane) containing hydrogen with a partial pressure of 1 MPa or more at the above temperature and pressure. CT test pieces (plate thickness 12.7 mm, plate width 25.4 mm) were used as test pieces, and were taken in a direction in which the machine notch introduction direction and the rolling direction of the steel material were parallel. Fatigue pre-cracks were introduced in the atmosphere under the conditions of frequency: 1 Hz, repeated load waveform: sine wave, control method: K value control, and stress ratio R: 0.1. After that, the test was carried out in hydrogen gas or a mixed atmosphere of hydrogen gas and natural gas. The fracture toughness test was carried out by the unloading-elastic compliance method using a single test piece. The crosshead displacement speed during loading was 0.002 mm/sec.
　評価結果を表３－１、３－２に示す。本発明例を満足する鋼材および鋼管は、全て水素誘起き裂進展下限界ＫＩＨが８０ＭＰａ・ｍ１／２以上である優れた水素中耐破壊靭性を示し、引張強さが５２０ＭＰａ以上を示した。なお、表３－１、３－２中の鋼管についても鋼材と同様の結果が得られている。 The evaluation results are shown in Tables 3-1 and 3-2. All of the steel materials and steel pipes satisfying the invention examples exhibited excellent fracture toughness in hydrogen, with a hydrogen-induced crack propagation threshold K IH of 80 MPa·m 1/2 or more, and a tensile strength of 520 MPa or more. The steel pipes in Tables 3-1 and 3-2 also showed similar results to the steel materials.
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000003
Figure JPOXMLDOC01-appb-T000003
Figure JPOXMLDOC01-appb-T000004
Figure JPOXMLDOC01-appb-T000004
Figure JPOXMLDOC01-appb-T000005
Figure JPOXMLDOC01-appb-T000005
Figure JPOXMLDOC01-appb-T000006
Figure JPOXMLDOC01-appb-T000006
　以下、本発明の効果を検証した実施例について、説明する。なお、以下の実施例において鋼材および鋼管を以下の製造条件で製造し、特性評価を行った。実施例１で用いた表１－１、１－２に示す鋼種Ｎｏ．２、４、８、１４、２２、３３を用いて、制御冷却工程まではそれぞれ実施例１（表２－１、２－２）で示す鋼管２、４、８、１４、２２、３３と同一の条件で製造し、鋼管成形も実施例１と同じ条件で実施し、脱水素処理条件を変化させたときの特性評価を行った。上記結果を表４に示す。 Below, examples verifying the effects of the present invention will be described. In the following examples, steel materials and steel pipes were manufactured under the following manufacturing conditions, and their characteristics were evaluated. Using steel types No. 2, 4, 8, 14, 22, and 33 shown in Tables 1-1 and 1-2 used in Example 1, they were manufactured under the same conditions as steel pipes 2, 4, 8, 14, 22, and 33 shown in Example 1 (Tables 2-1 and 2-2) up until the controlled cooling process, and steel pipe forming was also performed under the same conditions as in Example 1, and characteristics were evaluated when the dehydrogenation treatment conditions were changed. The results are shown in Table 4.
　実施例１で実施している鋼管Ｎｏ．２、４、８、１４、２２、３３の脱水素処理は、脱水素処理温度Ｔ（雰囲気温度）は表２－１、２－２に示す温度、時間で実施しており、これは、それぞれ表４の脱水素保持時間ｔがＹ、板厚中心温度Ｔｃにおける保持時間ｔｃがＮに相当する。 The dehydrogenation treatment of steel pipes No. 2, 4, 8, 14, 22, and 33 carried out in Example 1 was carried out at the dehydrogenation treatment temperature T (ambient temperature) and time shown in Tables 2-1 and 2-2, which correspond to the dehydrogenation holding time t in Table 4 being Y and the holding time tc at the plate thickness center temperature Tc being N, respectively.
　鋼管Ｎｏ．２Ａ、４Ａ、８Ａ、１４Ａ、２２Ａ、３３Ａは、脱水素処理温度Ｔはそれぞれ表４に示す温度とし、板厚中心温度Ｔｃが表４に示す脱水素処理温度Ｔに到達してからの保持時間ｔｃを（Ａ）式が満足するように実施した。 For steel pipes No. 2A, 4A, 8A, 14A, 22A, and 33A, the dehydrogenation temperature T was set to the temperature shown in Table 4, and the holding time tc after the center temperature Tc reached the dehydrogenation temperature T shown in Table 4 was set to satisfy formula (A).
　鋼管Ｎｏ．２Ｂ、４Ｂ、８Ｂ、１４Ｂ、２２Ｂ、３３Ｂは、脱水素処理温度Ｔはそれぞれ表４に示す温度であるが、雰囲気温度の保持時間ｔ、板厚中心温度Ｔｃが上記の脱水素処理温度Ｔに到達してからの保持時間ｔｃがともに上述している（Ａ）式を満足していない。 For steel pipes No. 2B, 4B, 8B, 14B, 22B, and 33B, the dehydrogenation temperature T is the temperature shown in Table 4, but the holding time t of the ambient temperature and the holding time tc after the temperature at the center of the plate thickness Tc reaches the above-mentioned dehydrogenation temperature T do not satisfy the above-mentioned formula (A).
　表４において、「脱水素保持時間ｔがＹ」は、脱水素処理温度Ｔ（雰囲気温度）は所定の温度とし、保持時間ｔが（Ａ）式を満足しており、「脱水素保持時間ｔがＮ」は、脱水素処理温度Ｔ（雰囲気温度）は所定の温度としているが、保持時間ｔが（Ａ）式を満足していない。また、「鋼材中心温度Ｔｃにおける保持時間ｔｃがＹ」は、板厚中心温度Ｔｃが所定の温度に到達してからの保持時間ｔｃが（Ａ）式を満足しており、「鋼材中心温度Ｔｃにおける保持時間ｔｃがＮ」は、板厚中心温度Ｔｃが所定の温度に到達するものの、Ｔｃが所定の温度に到達してからの保持時間ｔｃが（Ａ）式を満足していない。 In Table 4, "dehydrogenation holding time t is Y" means that the dehydrogenation temperature T (ambient temperature) is a predetermined temperature and the holding time t satisfies formula (A), and "dehydrogenation holding time t is N" means that the dehydrogenation temperature T (ambient temperature) is a predetermined temperature, but the holding time t does not satisfy formula (A). Also, "holding time tc at steel center temperature Tc is Y" means that the holding time tc after the plate thickness center temperature Tc reaches a predetermined temperature satisfies formula (A), and "holding time tc at steel center temperature Tc is N" means that the plate thickness center temperature Tc reaches a predetermined temperature, but the holding time tc after Tc reaches the predetermined temperature does not satisfy formula (A).
　水素中破壊靭性、引張強度の調査、組織、介在物の評価は、実施例１と同様の方法でおこなった。 Investigations of fracture toughness in hydrogen, tensile strength, and evaluation of structure and inclusions were performed in the same manner as in Example 1.
　本発明の発明例は、全て水素誘起き裂進展下限界ＫＩＨが８０ＭＰａ・ｍ１／２以上、引張強さが５２０ＭＰａ以上の条件を満足した。そのなかでも、脱水素処理条件がより好適な条件で実施される方が、水素中耐破壊靭性は優れていた。 All of the inventive examples of the present invention satisfied the conditions of a hydrogen-induced crack propagation threshold K IH of 80 MPa·m 1/2 or more and a tensile strength of 520 MPa or more. Among them, the examples in which the dehydrogenation treatment was performed under more suitable conditions had superior fracture toughness in hydrogen.
　なお、表４中の鋼管についても鋼材と同様の結果が得られている。 In addition, similar results were obtained for the steel pipes in Table 4 as for the steel materials.
Figure JPOXMLDOC01-appb-T000007
  
Figure JPOXMLDOC01-appb-T000007
  

Claims (8)

  1. 　質量％で、
    Ｃ：０．０２～０．１５％、
    Ｓｉ：０．０１～２．０％、
    Ｍｎ：０．５～１．５％、
    Ｐ：０．０００１～０．０１５％、
    Ｓ：０．０００２～０．００１５％、
    Ａｌ：０．００５～０．１５％、
    Ｏ：０．０１％以下、
    Ｎ：０．０１０％以下、
    Ｎｂ：０．１０％以下、
    Ｈ：０．０２ｐｐｍ以下を含み、
    あるいはさらに、
    Ｃａ：０～０．００５％、
    Ｎｉ：０～２．０％、
    Ｔｉ：０～０．１％、
    Ｃｕ：０～１．０％、
    Ｃｒ：０～１．０％、
    Ｍｏ：０～０．６０％、
    Ｗ：０～１．０％、
    Ｖ：０～０．１０％、
    Ｚｒ：０～０．０５０％、
    Ｍｇ：０～０．０１％、
    ＲＥＭ：０～０．０１％、
    Ｂ：０～０．００２０％、
    Ｔａ：０～０．２％、
    Ｈｆ：０～０．２％、
    Ｒｅ：０～０．００５％、
    Ｓｎ：０～０．３％、
    Ｓｂ：０～０．３％から選択される１種以上を含み、
    残部がＦｅおよび不可避的不純物元素である、化学組成を有し、
    ベイナイトおよびアスペクト比が２．０以上かつ長さが１０μｍ以上の介在物が１５個／１００ｍｍ２以下である金属組織を有し、
    鋼材表面から板厚中央の範囲における前記ベイナイトの最大粒径が２５μｍ以下であり、
    引張強度が５２０ＭＰａ以上であって、
    １ＭＰａ以上の高圧水素ガス環境において、水素誘起き裂進展下限界ＫＩＨが８０ＭＰａ・ｍ１／２以上である水素中破壊靭性に優れた高強度ラインパイプ用鋼材。     In mass percent,
    C: 0.02 to 0.15%,
    Si: 0.01 to 2.0%,
    Mn: 0.5 to 1.5%,
    P: 0.0001 to 0.015%,
    S: 0.0002 to 0.0015%,
    Al: 0.005 to 0.15%,
    O: 0.01% or less,
    N: 0.010% or less,
    Nb: 0.10% or less,
    H: 0.02 ppm or less,
    Or even more so:
    Ca: 0 to 0.005%,
    Ni: 0 to 2.0%,
    Ti: 0 to 0.1%,
    Cu: 0 to 1.0%,
    Cr: 0 to 1.0%,
    Mo: 0 to 0.60%,
    W: 0 to 1.0%,
    V: 0 to 0.10%,
    Zr: 0 to 0.050%,
    Mg: 0 to 0.01%,
    REM: 0 to 0.01%,
    B: 0 to 0.0020%,
    Ta: 0 to 0.2%,
    Hf: 0 to 0.2%,
    Re: 0 to 0.005%,
    Sn: 0 to 0.3%,
    Sb: one or more selected from 0 to 0.3%,
    The balance is Fe and unavoidable impurity elements,
    The metal structure has bainite and inclusions having an aspect ratio of 2.0 or more and a length of 10 μm or more at a ratio of 15 pieces/100 mm2 or less ,
    The maximum grain size of the bainite in the range from the steel surface to the center of the plate thickness is 25 μm or less,
    The tensile strength is 520 MPa or more,
    A high-strength linepipe steel material with excellent fracture toughness in hydrogen, having a hydrogen-induced crack propagation threshold K IH of 80 MPa·m 1/2 or more in a high-pressure hydrogen gas environment of 1 MPa or more.
  2. 　さらに、前記化学組成が、質量％で、
    Ｃａ：０．０００１～０．００５％、
    Ｎｉ：０．０１～２．０％、
    Ｔｉ：０．００５～０．１％、
    Ｃｕ：０．０１～１．０％、
    Ｃｒ：０．０１～１．０％、
    Ｍｏ：０．０１～０．６０％、
    Ｗ：０．０１～１．０％、
    Ｖ：０．０１～０．１０％、
    Ｚｒ：０．０００１～０．０５０％、
    Ｍｇ：０．０００１～０．０１％、
    ＲＥＭ：０．０００１～０．０１％、
    Ｂ：０．０００１～０．００２０％、
    Ｔａ：０．０００１～０．２％、
    Ｈｆ：０．０００１～０．２％、
    Ｒｅ：０．０００１～０．００５％、
    Ｓｎ：０．０００１～０．３％、
    Ｓｂ：０．０００１～０．３％である請求項１に記載の水素中破壊靭性に優れた高強度ラインパイプ用鋼材。     Further, the chemical composition comprises, in mass %,
    Ca: 0.0001 to 0.005%,
    Ni: 0.01 to 2.0%,
    Ti: 0.005 to 0.1%,
    Cu: 0.01 to 1.0%,
    Cr: 0.01 to 1.0%,
    Mo: 0.01 to 0.60%,
    W: 0.01 to 1.0%,
    V: 0.01 to 0.10%,
    Zr: 0.0001 to 0.050%,
    Mg: 0.0001 to 0.01%,
    REM: 0.0001 to 0.01%,
    B: 0.0001 to 0.0020%,
    Ta: 0.0001 to 0.2%,
    Hf: 0.0001 to 0.2%,
    Re: 0.0001 to 0.005%,
    Sn: 0.0001 to 0.3%,
    2. A steel material for high strength line pipes having excellent fracture toughness in hydrogen according to claim 1, wherein Sb is 0.0001 to 0.3%.
  3. 　残留オーステナイトが面積分率で０～３％であり、鋼材表面から板厚中央の範囲における前記ベイナイトが面積分率で９０％以上である請求項１または２に記載の水素中破壊靭性に優れた高強度ラインパイプ用鋼材。 A high-strength line pipe steel material with excellent fracture toughness in hydrogen according to claim 1 or 2, in which the area fraction of the retained austenite is 0-3%, and the area fraction of the bainite in the range from the steel surface to the center of the plate thickness is 90% or more.
  4. 　請求項１または２に記載の成分組成を有する鋳片を１０００～１２５０℃で加熱する加熱工程と、
    前記加熱工程で加熱された前記鋳片を、再結晶温度域での総圧下率が３５％以上５５％以下、かつ前記再結晶温度域での最終圧延パスの圧下率が１０％以上、かつ（再結晶温度－８０℃）以上における最終圧延パスの圧下率が１５％以上、さらに鋼板表面温度で圧延終了温度がＡｒ３変態点以上の条件で圧延する熱間圧延工程と、
    前記熱間圧延工程で得られた熱延鋼板を、冷却開始温度が前記熱延鋼板の表面温度でＡｒ３変態点以上、前記熱延鋼板の先端と尾端の冷却開始時間差が５０秒以内、７５０℃から５５０℃までの平均冷却速度が板厚中央温度で１５～５０℃／ｓ、冷却停止温度が２５０～６５０℃である条件で冷却する制御冷却工程と、
    を有する水素中破壊靭性に優れた高強度ラインパイプ用鋼材の製造方法。     A heating step of heating a slab having the component composition according to claim 1 or 2 at 1000 to 1250 ° C.;
    a hot rolling process in which the slab heated in the heating process is rolled under the conditions of a total rolling reduction in a recrystallization temperature range of 35% to 55%, a final rolling pass rolling reduction in the recrystallization temperature range of 10% or more, a final rolling pass rolling reduction at a temperature equal to or higher than (recrystallization temperature - 80 ° C.) of 15% or more, and a rolling end temperature at a steel plate surface temperature of the Ar3 transformation point or more;
    A controlled cooling process in which the hot-rolled steel sheet obtained in the hot rolling process is cooled under the conditions that the cooling start temperature is the Ar3 transformation point or higher at the surface temperature of the hot-rolled steel sheet, the cooling start time difference between the front end and the tail end of the hot-rolled steel sheet is within 50 seconds, the average cooling rate from 750 ° C. to 550 ° C. is 15 to 50 ° C./s at the plate thickness center temperature, and the cooling stop temperature is 250 to 650 ° C.;
    The present invention relates to a method for producing a high-strength steel material for line pipes having excellent fracture toughness in hydrogen.
  5. 　質量％で、
    Ｃ：０．０２～０．１５％、
    Ｓｉ：０．０１～２．０％、
    Ｍｎ：０．５～１．５％、
    Ｐ：０．０００１～０．０１５％、
    Ｓ：０．０００２～０．００１５％、
    Ａｌ：０．００５～０．１５％、
    Ｏ：０．０１％以下、
    Ｎ：０．０１０％以下、
    Ｎｂ：０．１０％以下、
    Ｈ：０．０２ｐｐｍ以下を含み、
    あるいはさらに、
    Ｃａ：０～０．００５％、
    Ｎｉ：０～２．０％、
    Ｔｉ：０～０．１％、
    Ｃｕ：０～１．０％、
    Ｃｒ：０～１．０％、
    Ｍｏ：０～０．６０％、
    Ｗ：０～１．０％、
    Ｖ：０～０．１０％、
    Ｚｒ：０～０．０５０％、
    Ｍｇ：０～０．０１％、
    ＲＥＭ：０～０．０１％、
    Ｂ：０～０．００２０％、
    Ｔａ：０～０．２％、
    Ｈｆ：０～０．２％、
    Ｒｅ：０～０．００５％、
    Ｓｎ：０～０．３％、
    Ｓｂ：０～０．３％から選択される１種以上を含み、
    残部がＦｅおよび不可避的不純物元素である、化学組成を有し、
    ベイナイトおよびアスペクト比が２．０以上かつ長さが１０μｍ以上の介在物が１５個／１００ｍｍ２以下である金属組織を有し、
    鋼管内面の表面から板厚中央の範囲における前記ベイナイトの最大粒径が２５μｍ以下であり、
    引張強度が５２０ＭＰａ以上であって、
    １ＭＰａ以上の高圧水素ガス環境において、水素誘起き裂進展下限界ＫＩＨが８０ＭＰａ・ｍ１/２以上である水素中破壊靭性に優れた高強度ラインパイプ用鋼管。     In mass percent,
    C: 0.02 to 0.15%,
    Si: 0.01 to 2.0%,
    Mn: 0.5 to 1.5%,
    P: 0.0001 to 0.015%,
    S: 0.0002 to 0.0015%,
    Al: 0.005 to 0.15%,
    O: 0.01% or less,
    N: 0.010% or less,
    Nb: 0.10% or less,
    H: 0.02 ppm or less,
    Or even more so:
    Ca: 0 to 0.005%,
    Ni: 0 to 2.0%,
    Ti: 0 to 0.1%,
    Cu: 0 to 1.0%,
    Cr: 0 to 1.0%,
    Mo: 0 to 0.60%,
    W: 0 to 1.0%,
    V: 0 to 0.10%,
    Zr: 0 to 0.050%,
    Mg: 0 to 0.01%,
    REM: 0 to 0.01%,
    B: 0 to 0.0020%,
    Ta: 0 to 0.2%,
    Hf: 0 to 0.2%,
    Re: 0 to 0.005%,
    Sn: 0 to 0.3%,
    Sb: one or more selected from 0 to 0.3%,
    The balance is Fe and unavoidable impurity elements,
    The metal structure has bainite and inclusions having an aspect ratio of 2.0 or more and a length of 10 μm or more at a ratio of 15 pieces/100 mm2 or less ,
    The maximum grain size of the bainite in the range from the surface of the steel pipe inner surface to the center of the plate thickness is 25 μm or less,
    The tensile strength is 520 MPa or more,
    A high-strength steel pipe for line pipes with excellent fracture toughness in hydrogen, having a hydrogen-induced crack propagation threshold K IH of 80 MPa·m 1/2 or more in a high-pressure hydrogen gas environment of 1 MPa or more.
  6. 　さらに、前記化学組成が、質量％で、
    Ｃａ：０．０００１～０．００５％、
    Ｎｉ：０．０１～２．０％、
    Ｔｉ：０．００５～０．１％、
    Ｃｕ：０．０１～１．０％、
    Ｃｒ：０．０１～１．０％、
    Ｍｏ：０．０１～０．６０％、
    Ｗ：０．０１～１．０％、
    Ｖ：０．０１～０．１０％、
    Ｚｒ：０．０００１～０．０５０％、
    Ｍｇ：０．０００１～０．０１％、
    ＲＥＭ：０．０００１～０．０１％、
    Ｂ：０．０００１～０．００２０％、
    Ｔａ：０．０００１～０．２％、
    Ｈｆ：０．０００１～０．２％、
    Ｒｅ：０．０００１～０．００５％、
    Ｓｎ：０．０００１～０．３％、
    Ｓｂ：０．０００１～０．３％である請求項５に記載の水素中破壊靭性に優れた高強度ラインパイプ用鋼管。     Further, the chemical composition comprises, in mass %,
    Ca: 0.0001 to 0.005%,
    Ni: 0.01 to 2.0%,
    Ti: 0.005 to 0.1%,
    Cu: 0.01 to 1.0%,
    Cr: 0.01 to 1.0%,
    Mo: 0.01 to 0.60%,
    W: 0.01 to 1.0%,
    V: 0.01 to 0.10%,
    Zr: 0.0001 to 0.050%,
    Mg: 0.0001 to 0.01%,
    REM: 0.0001 to 0.01%,
    B: 0.0001 to 0.0020%,
    Ta: 0.0001 to 0.2%,
    Hf: 0.0001 to 0.2%,
    Re: 0.0001 to 0.005%,
    Sn: 0.0001 to 0.3%,
    6. A steel pipe for high strength line pipe having excellent fracture toughness in hydrogen according to claim 5, wherein Sb is 0.0001 to 0.3%.
  7. 　高強度ラインパイプ用鋼管において、
    残留オーステナイトが面積分率で０～３％であり、鋼管内面の表面から板厚中央の範囲における前記ベイナイトが面積分率で９０％以上である請求項５または６に記載の水素中破壊靭性に優れた高強度ラインパイプ用鋼管。     In high-strength steel pipes for line pipes,
    7. A steel pipe for high strength line pipes having excellent fracture toughness in hydrogen according to claim 5 or 6, wherein the area fraction of retained austenite is 0 to 3%, and the area fraction of the bainite in the range from the surface of the inner surface of the steel pipe to the center of the plate thickness is 90% or more.
  8. 　請求項５または６に記載の成分組成を有する鋳片を１０００～１２５０℃で加熱する加熱工程と、
    前記加熱工程で加熱された前記鋳片を、再結晶温度域での総圧下率が３５％以上５５％以下、かつ前記再結晶温度域での最終圧延パスの圧下率が１０％以上、かつ（再結晶温度－８０℃）以上における最終圧延パスの圧下率が１５％以上、さらに鋼板表面温度で圧延終了温度がＡｒ３変態点以上の条件で圧延する熱間圧延工程と、
    該熱間圧延工程で得られた熱延鋼板を、冷却開始温度が前記熱延鋼板の表面温度でＡｒ３変態点以上、前記熱延鋼板の先端と尾端の冷却開始時間差が５０秒以内、７５０℃から５５０℃までの平均冷却速度が板厚中央温度で１５～５０℃／ｓ、冷却停止温度が２５０～６５０℃である条件で冷却する制御冷却工程と、
    該制御冷却工程後、前記熱延鋼板を曲げ加工し、両端部を突合せて溶接する造管工程、前記制御冷却工程後、前記熱延鋼板を冷間ロール成形により円筒状に成形し、該円筒状の周方向両端部を突合せて電縫溶接する造管工程のうちどちらか一方の造管工程と、
    を有する水素中破壊靭性に優れた高強度ラインパイプ用鋼管の製造方法。

          A heating step of heating a slab having the component composition according to claim 5 or 6 at 1000 to 1250 ° C.;
    a hot rolling process in which the slab heated in the heating process is rolled under the conditions of a total rolling reduction in a recrystallization temperature range of 35% to 55%, a final rolling pass rolling reduction in the recrystallization temperature range of 10% or more, a final rolling pass rolling reduction at a temperature equal to or higher than (recrystallization temperature - 80 ° C.) of 15% or more, and a rolling end temperature at a steel plate surface temperature of the Ar3 transformation point or more;
    A controlled cooling process in which the hot-rolled steel sheet obtained in the hot rolling process is cooled under the following conditions: a cooling start temperature is the Ar3 transformation point or higher at the surface temperature of the hot-rolled steel sheet, a cooling start time difference between the front end and the tail end of the hot-rolled steel sheet is within 50 seconds, an average cooling rate from 750°C to 550°C is 15 to 50°C/s at the center temperature of the sheet thickness, and a cooling stop temperature is 250 to 650°C;
    a pipe-making process in which, after the controlled cooling process, the hot-rolled steel sheet is bent and both ends are butted together and welded; or 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 and both circumferential ends of the cylindrical shape are butted together and electric resistance welded;
    A method for manufacturing high strength steel pipe for line pipe having excellent fracture toughness in hydrogen.
PCT/JP2023/035560 2022-09-29 2023-09-28 High-strength line pipe steel material having excellent fracture toughness in hydrogen, method for manufacturing same, steel tube for high-strength line pipes, and method for manufacturing same WO2024071358A1 (en)

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JP2012122103A (en) * 2010-12-09 2012-06-28 Sumitomo Metal Ind Ltd Thick steel plate excellent in hydrogen-induced crack resistance, brittle crack propagation arrest characteristic, and corrosion resistance
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JP2018012855A (en) * 2016-07-20 2018-01-25 新日鐵住金株式会社 Low alloy steel material, low alloy steel tube and container and method for producing the container
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WO2022030818A1 (en) * 2020-08-07 2022-02-10 주식회사 포스코 Steel material having excellent hydrogen embrittlement resistance and impact toughness and method for manufacturing same
JP2022068942A (en) * 2020-10-23 2022-05-11 Ｊｆｅスチール株式会社 Steel material for high-pressure hydrogen gas environment, and method for manufacturing the same
WO2022209896A1 (en) * 2021-03-30 2022-10-06 Ｊｆｅスチール株式会社 Steel pipe for high-pressure hydrogen, container for high-pressure hydrogen, and method for manufacturing said steel pipe

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JP2012122103A (en) * 2010-12-09 2012-06-28 Sumitomo Metal Ind Ltd Thick steel plate excellent in hydrogen-induced crack resistance, brittle crack propagation arrest characteristic, and corrosion resistance
WO2014156187A1 (en) * 2013-03-29 2014-10-02 Ｊｆｅスチール株式会社 Steel material and hydrogen container as well as manufacturing methods therefor
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