JP7838630B2 - Steel materials for line pipes and their manufacturing method, steel pipes for line pipes and their manufacturing method - Google Patents

Description

This invention relates to a steel material for line pipes suitable for applications such as line pipes for transporting hydrogen gas, a method for manufacturing the same, a steel pipe for line pipes, and a method for manufacturing the same.

Existing energy infrastructure includes pipelines for transporting natural gas. These steel materials have been required to suppress hydrogen-induced cracking in sour environments. Meanwhile, in recent years, hydrogen has attracted significant global attention as a clean energy source for building a decarbonized society. Therefore, to transport large quantities of hydrogen gas, the construction of hydrogen gas transport networks is being considered, either by mixing natural gas with some hydrogen into the pipelines or by using hydrogen gas as a substitute for pressurized transport. The transport pressure during operation of these pipelines is expected to be high, ranging from 1 to 40 MPa, meaning the pipelines will be exposed to a high-pressure hydrogen gas environment. In such environments, there is a concern that steel materials may undergo "hydrogen embrittlement," where hydrogen penetrates the steel, degrading its properties. Therefore, in addition to the high toughness and sour resistance required for conventional pipelines, steel materials need to possess the resistance to hydrogen embrittlement required in a hydrogen gas environment.

For steel structures used in high-pressure hydrogen gas environments, austenitic stainless steels such as SUS316L have traditionally been used because they are less susceptible to hydrogen embrittlement than low-alloy steels. However, austenitic stainless steels such as SUS316L are expensive and have low strength. Therefore, designing them to withstand high hydrogen pressure requires thicker walls, resulting in higher prices for the hydrogen structures themselves. For this reason, there has been a strong demand for low-alloy steel materials that are more cost-effective and can withstand high-pressure hydrogen gas environments for hydrogen steel structures.

In response to such demands, for example, the steel for high-pressure hydrogen environments described in Patent Document 1 is a steel used in a high-pressure hydrogen environment, and by setting the Ca/S ratio to less than 1.5 or 11 or more, the diffusible hydrogen concentration ratio is reduced and embrittlement due to diffusible hydrogen is suppressed.

Furthermore, Patent Document 2 describes a technology that, by using low-alloy high-strength steel adjusted to a specific component composition, demonstrates that in the tensile strength range of 900 to 950 MPa in air, the reduction of area and elongation values in a 45 MPa hydrogen atmosphere are greater than those of JIS G3128SHY685NS, and that it exhibits superior resistance to embrittlement in high-pressure hydrogen environments.

Furthermore, Patent Document 3 describes a Cr-Mo-based high-strength low-alloy steel that, by tempering it at a relatively high temperature of 560 to 580°C, and adjusting the grain size after tempering to a grain size of 8.4 or higher, and the tensile strength to an extremely narrow range of 900 to 950 MPa, exhibits excellent elongation and reduction characteristics even in a 45 MPa hydrogen atmosphere, resulting in a low-alloy high-strength steel with excellent resistance to high-pressure hydrogen environment embrittlement.

Furthermore, the low-alloy steel for high-pressure hydrogen gas environments proposed in Patent Document 4 involves adding V, increasing the Mo content compared to existing steels, and raising the tempering temperature to utilize V-Mo carbides. This improves the carbide morphology at grain boundaries, significantly enhancing resistance to hydrogen environment embrittlement.

Furthermore, Patent Document 5 proposes a steel for high-pressure hydrogen gas storage containers with excellent hydrogen resistance. According to the technology described in Patent Document 5, by performing long-term stress-relieving annealing after normalizing during steel plate manufacturing, MC-type carbides (Mo,V)C are dispersed and precipitated in a fine and high-density manner, thereby improving the hydrogen resistance of the steel, such as its resistance to hydrogen embrittlement.

Furthermore, Patent Document 6 proposes a steel material in which the metallic structure is predominantly bainite with an area fraction of 90% or more, and cementite with an average grain size of 50 nm or less and an average aspect ratio of 3 or less is dispersed and precipitated within the bainite.

Furthermore, Non-Patent Document 1 contains the fatigue strength values for low-alloy steel.

Japanese Patent Publication No. 2005-2386 Japanese Patent Publication No. 2009-46737 Japanese Patent Publication No. 2009-275249 Japanese Patent Publication No. 2009-74122 Japanese Patent Publication No. 2010-37655 Japanese Patent Publication No. 2012-107332

Matsunaga et al. , Int J Hydrogen Energy, Vol. 40 (2015), p. 5739-5748 Omura, Hydrogenation Method for Reproducing Hydrogen Intrusion in Real Environments, The Japan Institute of Metals, Vol. 84, No. 9, (2020), pp. 285-294. Japan Heat Treatment Technology Association (author), Introduction to the Structure and Properties of Metallic Materials - Heat Treatment and Microstructure Control for Maximizing Material Potential, 2004. Michiya Okada et al., Hydrogen permeation in iron measured electrochemically, Journal of the Japan Institute of Metals, Vol. 50, No. 2 (1986), pp. 201-207.

The pressure inside a line pipe fluctuates during operation and undergoes periodic shutdowns, subjecting the structure to repeated stress. Therefore, when designing steel structures such as line pipes, it is essential to consider fatigue failure. However, as shown in Non-Patent Document 1, it is known that the fatigue life of materials decreases under high-pressure hydrogen environments. This means that if line pipe materials are designed based on conventional natural gas line pipes, the service life of the line pipe material will be reduced. However, while the above-mentioned conventional technology can suppress hydrogen-induced cracking in sour environments, it cannot sufficiently increase the fatigue life in hydrogen gas. In other words, there is a problem in achieving both suppression of hydrogen-induced cracking in sour environments and high fatigue strength in hydrogen gas.

In view of the problems of the prior art described above, the present invention aims to provide a steel material for line pipes and a method for manufacturing the same, as well as a steel pipe for line pipes and a method for manufacturing the same, which are suitable for use in steel structures used in high-pressure hydrogen gas environments, such as line pipes for 100% hydrogen gas or natural gas containing hydrogen with a hydrogen partial pressure of 1 MPa or more (natural gas is a gas whose main components are hydrocarbons such as methane and ethane), and which are high in strength and have excellent resistance to hydrogen embrittlement in high-pressure hydrogen gas environments, in view of the problems of the prior art described above. The hydrogen environment is assumed to be a high-pressure hydrogen gas environment of 1 MPa or more or a natural gas environment containing hydrogen with a hydrogen partial pressure of 1 MPa or more (the main components are hydrocarbons such as methane and ethane).

In this context, "excellent resistance to hydrogen embrittlement under high-pressure hydrogen gas conditions" means that the fatigue crack propagation rate da/dN mm/cycle, obtained by conducting fatigue tests in accordance with ASTM E647, Fatigue Testing, under both conditions of room temperature (20 ± 10°C), hydrogen gas pressure of 1 MPa or higher, or a mixed atmosphere of natural gas (mainly composed of hydrocarbons such as methane and ethane) containing hydrogen as a partial pressure of 1 MPa or higher, with a frequency of 1 Hz, repetition waveform: sine wave, control method: load control, load condition: uniaxial tension, and stress ratio: R = 0.1, is 2.0 × ^10⁻³ mm/cycle or less at ΔK = 25 MPa.

Furthermore, under the above conditions, if the fatigue crack propagation rate da/dN mm/cycle is 2.0 × ^10⁻³ mm/cycle or less, it becomes possible to design long-life steel structures for hydrogen, such as line pipes, within the plate thickness range that can be manufactured using processes for manufacturing steel pipes such as seamless steel pipes and UOE.

Furthermore, the term "steel materials" as used here includes thin steel plates, thick steel plates, seamless steel pipes, electric resistance welded steel pipes, structural steel, steel bars, etc.

The inventors diligently researched the conditions that steel materials must satisfy to obtain line pipe steel materials and line pipe steel tubes with excellent hydrogen embrittlement resistance. As a result, they found that the fatigue crack propagation rate is greatly influenced by the accumulation of hydrogen at the crack tip and the stress at the crack tip (stress intensity factor), and that reducing the hydrogen solid solubility in the steel to 0.05 ppm/√P or less significantly reduces the fatigue crack propagation rate in hydrogen. Furthermore, the fatigue crack propagation rate in hydrogen accelerates as the amount of hydrogen accumulated at the crack tip increases. And the smaller the hydrogen diffusion coefficient, the greater the amount of hydrogen accumulated at the crack tip, and the greater the fatigue crack propagation rate. The inventors analyzed the relationship between fatigue crack propagation rate and hydrogen diffusion coefficient in detail and found that when the hydrogen diffusion coefficient at room temperature is smaller than 1.5 × ^{10⁻¹⁰ m²} /s, the fatigue crack propagation rate in hydrogen increases significantly. Based on the above findings, the inventors have invented new high-strength line pipe steel materials and line pipe steel tubes. Furthermore, the steel material and steel pipe of the present invention have high strength, and in the present invention, high strength refers to a tensile strength of 520 MPa or more.

The gist of this invention is as follows:
[1] In mass percent,
C: 0.02-0.15%,
Si: 0.01-2.0%,
Mn: 0.5-1.8%,
P: 0.0001-0.015%,
S: 0.0002-0.0015%,
Al: 0.005-0.15%,
O: 0.01% or less,
N: 0.010% or less,
H: Includes 0.02 ppm or less,
Or, furthermore,
Nb: 0 to 0.10%,
Ca: 0-0.005%,
Ni: 0-2.0%,
Ti: 0 to 0.1%,
Cu: 0 to 1.0%,
Cr: 0-1.0%,
Mo: 0 to 0.60%,
W: 0-1.0%,
V: 0-0.10%,
Zr: 0 to 0.050%,
REM: 0-0.01%,
Mg: 0 to 0.01%,
B: 0 to 0.0020%,
Hf: 0-0.2%,
Ta: 0-0.2%,
Re: 0 to 0.005%,
Sn: 0-0.3%,
Sb: Contains one or more selected from 0 to 0.3%,
It has a chemical composition in which the remainder is Fe and unavoidable impurity elements.
A steel material for line pipes having a retained austenite content of 0-3% by area fraction, a hydrogen diffusion coefficient of 1.5 × ^{10⁻¹⁰ m²} /s or higher at room temperature, and a hydrogen solid solubility of 0.05 mass ppm/√P or less.
[2] Furthermore, the chemical composition is, in mass%,
Nb: 0.001 to 0.10%,
Ca: 0.0001-0.005%,
Ni: 0.01-2.0%,
Ti: 0.005-0.1%,
Cu: 0.01 to 1.0%,
Cr: 0.01-1.0%,
Mo: 0.01 to 0.60%,
W: 0.01-1.0%,
V: 0.01-0.10%,
Zr: 0.0001 to 0.050%,
REM: 0.0001-0.01%,
Mg: 0.0001-0.01%,
B: 0.0001 to 0.0020%,
Hf: 0.0001-0.2%,
Ta: 0.0001-0.2%,
Re: 0.0001-0.005%,
Sn: 0.0001-0.3%,
The steel material for line pipes described in [1], wherein Sb: 0.0001 to 0.3%.
[3] The steel material for line pipes according to [1] or [2], wherein the material has bainite or martensite at the 1/4 thickness position, and the bainite accounts for 90% or more of the area fraction, or the martensite accounts for 90% or more of the area fraction.
[4] A heating step of heating a steel material having the chemical composition described in [1] or [2] above to 1000 to 1250°C,
A hot rolling process is performed in which the steel material heated in the above heating step is rolled at a rolling completion temperature of Ar ₃ or higher.
A controlled cooling process is performed to cool the hot-rolled steel sheet obtained in the hot-rolling process under the following conditions: the cooling start temperature is Ar ₃ or higher at the steel sheet surface temperature, the difference in cooling start time between the leading and trailing ends 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 temperatures 0.25 mm below the steel sheet surface and in the center of the sheet thickness, and the cooling stop temperature is 250 to 650°C.
A stabilization process for stabilizing the steel sheet obtained in the controlled cooling process, or a dehydrogenation process for dehydrogenating the steel sheet obtained in the controlled cooling process,
A method for manufacturing steel materials for line pipes.
[5] In steel pipes for line pipes,
In mass percent,
C: 0.02-0.15%,
Si: 0.01-2.0%,
Mn: 0.5-1.8%,
P: 0.0001-0.015%,
S: 0.0002-0.0015%,
Al: 0.005-0.15%,
O: 0.01% or less,
N: 0.010% or less,
H: Includes 0.02 ppm or less,
Or, furthermore,
Nb: 0 to 0.10%,
Ca: 0-0.005%,
Ni: 0-2.0%,
Ti: 0 to 0.1%,
Cu: 0 to 1.0%,
Cr: 0-1.0%,
Mo: 0 to 0.60%,
W: 0-1.0%,
V: 0-0.10%,
Zr: 0 to 0.050%,
REM: 0-0.01%,
Mg: 0 to 0.01%,
B: 0 to 0.0020%,
Hf: 0-0.2%,
Ta: 0-0.2%,
Re: 0 to 0.005%,
Sn: 0-0.3%,
Sb: Contains one or more selected from 0 to 0.3%,
It has a chemical composition in which the remainder is Fe and unavoidable impurity elements.
A steel pipe for line pipes having a retained austenite content of 0-3% by area fraction, a hydrogen diffusion coefficient of 1.5 × ^{10⁻¹⁰ m²} /s or higher at room temperature, and a hydrogen solid solubility of 0.05 mass ppm/√P or less.
[6] Furthermore, the chemical composition is, in mass%,
Nb: 0.001 to 0.10%,
Ca: 0.0001-0.005%,
Ni: 0.01-2.0%,
Ti: 0.005-0.1%,
Cu: 0.01 to 1.0%,
Cr: 0.01-1.0%,
Mo: 0.01 to 0.60%,
W: 0.01-1.0%,
V: 0.01-0.10%,
Zr: 0.0001 to 0.050%,
REM: 0.0001-0.01%,
Mg: 0.0001-0.01%,
B: 0.0001 to 0.0020%,
Hf: 0.0001-0.2%,
Ta: 0.0001-0.2%,
Re: 0.0001-0.005%,
Sn: 0.0001-0.3%,
Steel pipe for line pipes as described in [5], wherein Sb: 0.0001 to 0.3%.
[7] A steel pipe for line pipes according to [5] or [6], wherein the steel pipe has bainite or martensite at a position 1/4 of the wall thickness from the inner surface of the steel pipe, and the area fraction of bainite is 90% or more or the area fraction of martensite is 90% or more.
[8] A heating step of heating a steel material having the chemical composition described in [5] or [6] to 1000 to 1250°C,
A hot rolling process is performed in which the steel material heated in the above heating step is rolled at a rolling completion temperature of Ar ₃ or higher.
A controlled cooling process is performed to cool the hot-rolled steel sheet obtained in the hot-rolling process under the following conditions: the cooling start temperature is Ar ₃ or higher at the steel sheet surface temperature, the difference in cooling start time between the leading and trailing ends 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 temperatures 0.25 mm below the steel sheet surface and in the center of the sheet thickness, and the cooling stop temperature is 250 to 650°C.
After the controlled cooling process, a pipe-making process is performed in which the hot-rolled steel sheet is bent and both ends are butt-welded; or a pipe-making process is performed in which the hot-rolled steel sheet is formed into a cylindrical shape by cold roll forming and both ends of the cylindrical shape are butt-welded using electric resistance welding;
Either a stabilization process for stabilizing the steel pipes obtained in the pipe manufacturing process, or a dehydrogenation process for dehydrogenating the steel pipes obtained in the pipe manufacturing process,
A method for manufacturing steel pipes for line pipes.

According to the present invention, steel materials with significantly improved hydrogen embrittlement resistance under high-pressure hydrogen gas environments can be easily and simply manufactured, yielding remarkable industrial benefits. Furthermore, the present invention significantly improves the hydrogen embrittlement resistance of steel structures such as high-pressure hydrogen gas line pipes, improving fatigue resistance and greatly contributing to extending the lifespan of steel structures.

Next, a method for carrying out the present invention will be specifically described. The following description illustrates preferred embodiments of the present invention, and the present invention is not limited in any way by this description. A steel material will be specifically described as the first embodiment, followed by a UOE steel pipe, an example of a steel pipe of the present invention, as the second embodiment, and an electric resistance welded steel pipe, an example of a steel pipe of the present invention, as the third embodiment.

First Embodiment [Component Composition]
The reasons for the limitations on the chemical composition of the steel material of the present invention are explained below. In the following explanation, "%" refers to "mass percent" unless otherwise specified.

C: 0.02-0.15%
While carbon (C) effectively contributes to improving strength, sufficient strength cannot be ensured if the C content is less than 0.02%. Therefore, the C content should be 0.02% or more. Preferably, the C content is 0.03% or more. On the other hand, if it exceeds 0.15%, weldability decreases. Therefore, the C content should be limited to 0.15% or less. Preferably, the C content is 0.13% or less. Furthermore, if it exceeds 0.08%, the hardness of the surface layer and central segregation increases during accelerated cooling, which may lead to deterioration of SSCC resistance and HIC resistance. Toughness may also deteriorate. Therefore, more preferably, the C content is 0.08% or less. Even more preferably, the C content is 0.05% or less.

Si: 0.01~2.0%
Si is added for deoxidation, but if the content is less than 0.01%, the deoxidation effect is insufficient. Therefore, the Si content is 0.01% or more. Preferably, the Si content is 0.08% or more, and more preferably 0.10% or more. On the other hand, if it exceeds 2.0%, the effect saturates, so the Si content is 2.0% or less. Preferably, the Si content is 1.8% or less, and more preferably 1.0% or less. Furthermore, if it exceeds 0.50%, the toughness and weldability deteriorate, and the hydrogen solid solubility increases, so the Si content is even more preferably 0.50% or less.

Mn: 0.5-1.8%
While manganese (Mn) effectively contributes to improving strength and toughness, its additive effect is poor at concentrations below 0.5%. Therefore, the Mn content should be 0.5% or higher. Preferably, the Mn content is 0.6% or higher, more preferably 0.7% or higher, and even more preferably 0.8% or higher. On the other hand, if the Mn content exceeds 1.8%, the hardness of the surface layer and central segregation increases during controlled cooling, resulting in deterioration of resistance to SSCC (sulfide stress corrosion cracking), HIC (hydrogen-induced cracking), and hydrogen embrittlement. Weldability also deteriorates, and hydrogen solid solubility increases. Therefore, the Mn content should be limited to 1.8% or less. Preferably, the Mn content is 1.5% or less, more preferably 1.4% or less, and even more preferably 1.3% or less.

P:0.0001~0.015%
P is an unavoidable impurity element that degrades weldability and increases the hardness of the central segregation zone, thereby increasing HIC resistance and hydrogen solid solubility, and thus degrading hydrogen embrittlement resistance. This tendency becomes significant above 0.015%, so the upper limit for P content is set at 0.015%. A P content of 0.010% or less is preferable, and more preferably, a P content of 0.008% or less. A lower content is better, but from the viewpoint of refining costs, the P content should be 0.0001% or more.

S: 0.0002-0.0015%
S is an unavoidable impurity element, and in steel, it becomes a MnS inclusion, increasing HIC resistance and hydrogen solid solubility, thereby degrading hydrogen embrittlement resistance. Therefore, a low amount is preferable, but up to 0.0015% is acceptable. For this reason, the S content should be 0.0015% or less. A S content of 0.0010% or less is preferable, and 0.0008% or less is more preferable. A lower content is better, but from the viewpoint of refining costs, it should be 0.0002% or more.

Al: 0.005-0.15%
Al is added as a deoxidizing agent, but since there is no effect if the amount is less than 0.005%, the Al content should be 0.005% or more. The Al content is preferably 0.010% or more, and more preferably 0.030% 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 should be limited to 0.15% or less. The Al content is preferably 0.10% or less, more preferably 0.08% or less, and even more preferably 0.05% or less.

O: 0.01% or less. While it is preferable to have less oxygen because it causes the formation of oxide inclusions, an oxygen content of 0.01% or less is not problematic. Therefore, the oxygen content is set to 0.01% or less. Preferably, the oxygen content is 0.005% or less. More preferably, the oxygen content is less than 0.003%. There is no particular lower limit, but since reducing oxygen to 0% increases costs, 0.001% or more is preferred.

N: 0.010% or less The effect of N on the fatigue properties of steel is small, and if the N content is 0.010% or less, the effects of the present invention are not impaired from the viewpoint of toughness. Therefore, the N content is set to 0.010% or less. It is preferable that the N content be 0.008% or less, and more preferably 0.006% or less. It is even more preferable that the N content be 0.004% or less. On the other hand, from the viewpoint of improving toughness, a low N content is desirable, but excessive reduction increases steelmaking costs, so it is preferable that the N content be 0.001% or more.

H: 0.02 ppm or less. H may be introduced into the steel material during various processes in manufacturing. If the amount introduced is high, the risk of crack formation after solidification increases and fatigue crack propagation may be accelerated. These effects are not a problem if the H content is 0.02 ppm or less, so the H content should be 0.02 ppm or less. Preferably, the H content is 0.01 ppm or less. More preferably, it is 0.005 ppm or less. Even more preferably, the H content is 0.003 ppm or less. There is no particular lower limit, but since a value of less than 0.001 ppm will increase costs, an H content of 0.001 ppm or more is preferred.
The hydrogen content refers to the residual hydrogen content after molding of steel materials, steel pipes, UOE, etc.

The component composition of this disclosure may optionally contain one or more elements selected from Nb, Ca, Ni, Ti, Cu, Cr, Mo, W, V, Zr, REM, Mg, B, Hf, Ta, Re, Sn, and Sb within the following ranges in order to further improve the strength and toughness of the steel sheet.

Nb: 0-0.10%
Nb is an element that contributes to increasing the strength of steel, but if the content exceeds 0.10%, the effect saturates and becomes a factor in cost increase. Therefore, when Nb is included, the Nb content should be 0.10% or less. Preferably, the Nb content should be 0.08% or less. More preferably, the Nb content should be 0.06% or less. For cost reduction, it is even more preferable that the Nb content be 0.05% or less. When Nb is included, the Nb content may be 0% or more, but in order to obtain the above effect, it is preferable that the content be 0.001% or more. More preferably, the Nb content should be 0.01% or more.

Ca: 0-0.005%
Ca is an effective element for improving HIC resistance by controlling the morphology of sulfide inclusions. However, if the amount exceeds 0.005%, the effect not only saturates but also deteriorates HIC resistance due to a decrease in the cleanliness of the steel. Therefore, when Ca is included, the amount of Ca should be 0.005% or less. A Ca content of 0.003% or less is preferable. A Ca content of 0.002% or less is more preferable. When Ca is included, the Ca content may be 0% or more, but if it is less than 0.0001%, the effect of its inclusion is insufficient, so a Ca content of 0.0001% or more is preferable. A Ca content of 0.001% or more is more preferable.

Ni: 0-2.0%
Ni is an effective element for improving toughness and increasing strength, but if it is present in amounts exceeding 2.0%, it tends to generate fine cracks called Fischer cracks in environments with low hydrogen sulfide partial pressures of less than 1 bar. Therefore, when Ni is included, the Ni content should be 2.0% or less. A Ni content of 1.5% or less is preferable, 1.2% or less is more preferable, and 1.0% or less is even preferable. A Ni content of 0.1% or less is preferable. Most preferably, it should be 0.02% or less. When Ni is included, the Ni content may be 0% or more, but in order to obtain the above effects, it is preferable to have a Ni content of 0.01% or more.

Ti: 0-0.1%
Ti contributes to increasing the strength of steel, but its effect saturates when its content exceeds 0.1%, leading to increased costs. Therefore, when Ti is included, the Ti content should be 0.1% or less. To control costs, it is even more preferable to have a Ti content of 0.05% or less. When Ti is included, the Ti content may be 0% or more, but to obtain the above-mentioned effects, it is preferable to have a Ti content of 0.005% or more. A Ti content of 0.008% or more is even more preferable.

Cu: 0-1.0%
Cu is an effective element for improving toughness and increasing strength, but if the content is too high, weldability and hydrogen solid solubility increase, which deteriorates hydrogen embrittlement resistance. Therefore, when Cu is included, the Cu content should be 1.0% or less. A Cu content of 0.5% or less is preferable. A Cu content of 0.3% or less is more preferable, and 0.2% or less is even preferable. When Cu is included, the Cu content may be 0% or more, but to obtain this effect, it is preferable to have a Cu content of 0.01% or more.

Cr: 0-1.0%
Like manganese, chromium (Cr) is an effective element for obtaining sufficient strength even at low carbon content. However, if the content is too high, the hardenability becomes excessive, leading to increased resistance to SSCC (Steel Sulfur Carbon) and hydrogen solid solubility, which deteriorates hydrogen embrittlement resistance. Weldability also deteriorates. Therefore, when Cr is included, the Cr content should be 1.0% or less. A Cr content of 0.8% or less is preferable, more preferably 0.5% or less, and even more preferably 0.1% or less. When Cr is included, the Cr content may be 0% or more, but to obtain this effect, it is preferable to have a Cr content of 0.01% or more, and more preferably 0.05% or more.

Mo: 0~0.60%
Mo is an effective element for improving toughness and increasing strength, and is effective for improving resistance to SSCC regardless of hydrogen sulfide partial pressure. However, if the content is too high, the hardenability becomes excessive, and the resistance to hydrogen embrittlement deteriorates due to the increased resistance to SSCC and hydrogen solid solubility. Weldability also deteriorates. For this reason, the Mo content should be 0.60% or less. Preferably, the Mo content should be 0.50% or less, and more preferably 0.40% or less. The Mo content may be 0% or more, but to obtain the above effects, it is preferable to contain 0.01% or more Mo, and more preferably 0.10% or more.

W: 0-1.0%
While W contributes to increasing the strength of steel, the effect saturates when the W content exceeds 1.0%, leading to increased costs. Therefore, when W is included, the W content should be 1.0% or less. Preferably, the W content should be 0.8% or less. To control costs, it is preferable to have a W content of 0.5% or less. The W content may be 0% or more, but to obtain the above-mentioned effects, when W is included, it is preferable to have a W content of 0.01% or more.

V: 0-0.10%
V is an element that can be optionally included to increase the strength and toughness of steel. If the amount exceeds 0.10%, the toughness of the weld and the hydrogen solid solubility increase, which deteriorates the resistance to hydrogen embrittlement, so if V is included, it should be 0.10% or less. It is preferable that the V content be 0.08% or less. It is more preferable that the V content be 0.06% or less, and even more preferable that it be 0.03% or less. The V content may be 0% or more, but if the content is less than 0.01%, the effect of including it is poor, so if V is included, it is preferable that it be 0.01% or more.

Zr: 0-0.050%, REM: 0-0.01%, Mg: 0-0.01%
Zr, REM, and Mg are elements that can be optionally included to enhance toughness through grain refinement or to improve crack resistance through control of inclusion properties. Since the effects of these elements saturate when Zr exceeds 0.050% and REM and Mg exceed 0.01%, when these elements are included, the Zr content should be 0.050% or less, and the REM and Mg content should be 0.01% or less. Specifically, when these elements are included, the Zr content should be 0.050% or less. A Zr content of 0.0040% or less is preferable. A Zr content of 0.0030% or less is more preferable. Furthermore, when REM is included, the REM content should be 0.01% or less. A REM content of 0.0040% or less is preferable. A REM content of 0.0030% or less is more preferable. Furthermore, when Mg is included, the Mg content should be 0.01% or less. A Mg content of 0.0040% or less is preferable. It is more preferable that the Mg content be 0.0030% or less. The content of these elements may be 0% or more, but since the effect of their inclusion is poor if the content is less than 0.0001%, it is preferable that it be 0.0001% or more. In other words, it is preferable that the Zr content be 0.0001% or more. It is more preferable that the Zr content be 0.0005% or more. Also, it is preferable that the REM content be 0.0001% or more. It is more preferable that the REM content be 0.0005% or more. It is preferable that the Mg content be 0.0001% or more. It is more preferable that the Mg content be 0.0005% or more.

B: 0-0.0020%
B is an element that improves hardenability, contributing to increased strength of steel pipes, suppressing coarsening of prior austenite grains, and improving various properties of the material. On the other hand, if the B content exceeds 0.0020%, the effect saturates and becomes a factor in cost increase, so if B is included, the B content should be 0.0020% or less. Preferably, the B content should be 0.0015% or less. More preferably, the B content should be 0.0012% or less. For cost reduction, it is even more preferable to have a B content of 0.0010% or less. The B content may be 0% or more, but in order to obtain the above effects, if B is included, it is preferable to have a content of 0.0001% or more. More preferably, it is 0.0005% or more.

Hf: 0-0.2%, Ta: 0-0.2%
These elements contribute to increasing the strength of steel. Since the effect saturates when the content exceeds 0.2%, leading to increased costs, the content of these elements should be 0.2% or less when included. Specifically, if included, Hf should be 0.2% or less. Preferably, Hf should be 0.1% or less. More preferably, Hf should be 0.05% or less. Also, if included, Ta should be 0.2% or less. Preferably, Ta should be 0.1% or less. More preferably, Ta should be 0.05% or less. For cost reduction, the content should be 0.01% or less. While the content of these elements may be 0% or more, to obtain the aforementioned effects, if included, the Hf content is preferably 0.0001% or more. More preferably, the Hf content is 0.001% or more. Also, the Ta content is preferably 0.0001% or more. More preferably, the Ta content is 0.001% or more.

Re: 0~0.005%
Re contributes to increasing the strength of steel materials, but if the content exceeds 0.005%, the effect saturates and becomes a factor in cost increase; therefore, if it is included, it should be 0.005% or less. Preferably, the Re content should be 0.003% or less. More preferably, the Re content should be 0.002% or less. The Re content may be 0% or more, but in order to obtain the above effect, if it is included, the content should be 0.0001% or more. Preferably, it is 0.001% or more.

Sn: 0 to 0.3%, Sb: 0 to 0.3%,
These elements contribute to increasing the strength and hardenability of steel, but if their content exceeds 0.3%, the effect saturates, leading to increased costs. Therefore, if they are included, the content should be 0.3% or less. Specifically, the Sn content should be 0.3% or less. Preferably, the Sn content should be 0.2% or less. More preferably, the Sn content should be 0.1% or less. For cost reduction, it is even more preferable to have a content of 0.01% or less. The Sb content should also be 0.3% or less. Preferably, the Sb content should be 0.2% or less. More preferably, the Sb content should be 0.1% or less. For cost reduction, it is even more preferable to have a Sb content of 0.01% or less. The Sn and Sb content may be 0% or more, but to obtain the above effects, if they are included, it is preferable that the Sn content be 0.0001% or more. More preferably, the Sn content is 0.001% or more. Also, it is preferable that the Sb content be 0.0001% or more. More preferably, the Sb content is 0.0010% or more.

In the composition of steel plates and steel pipes, the remainder of the components (elements) other than those mentioned above consists of Fe and unavoidable impurity elements.

The following describes the microstructure of the steel material of the present invention.

Metallic structure: Retained austenite present in area fractions of 0-3%
The presence of retained austenite in steel can increase the hydrogen content and thus increase susceptibility to hydrogen embrittlement. Furthermore, if retained austenite transforms into martensite due to stress loading during use, the martensite is very hard and therefore prone to hydrogen cracking, potentially causing cracks to initiate from the martensite portion. In this invention, by limiting the retained austenite to 3% or less in area fraction, the fatigue crack propagation rate can be reduced, leading to improved resistance to hydrogen embrittlement. For this reason, the retained austenite is limited to 3% or less. Preferably, it is 2% or less. More preferably, it is 1% or less. The retained austenite may be 0%.

At a position where the steel plate thickness is 1/4 of the way through (or at a position where the wall thickness from the inner surface of the steel pipe is 1/4 of the way through in the case of a steel pipe), bainite or martensite is present, wherein the area fraction of bainite is 90% or more, or the area fraction of martensite is 90% or more (preferred).
To achieve high strength with a tensile strength of 520 MPa or more, the steel microstructure is preferably bainite or martensite. On the other hand, when soft and hard phases are mixed in the steel material, fatigue damage preferentially accumulates in the soft phase, making crack initiation more likely and reducing the fatigue limit stress. In a hydrogen environment, local deformation is promoted, so fatigue damage to the soft phase is further accelerated, and the hydrogen embrittlement resistance in hydrogen is further reduced. As a result, it becomes difficult to achieve a hydrogen fatigue crack propagation rate da/dNmm/cycle of 2.0 × ^10⁻³ mm/cycle or less at ΔK = 25 MPa. To improve this, it is necessary to reduce the relative proportion of the soft phase. For this reason, the metallic microstructure is preferably a single structure of bainite or martensite, and preferably has either bainite or martensite, with the area fraction of that structure being 90% or more. It is more preferable that the area fraction of the structure be 92% or more, and even more preferable that it be 95% or more. There is no particular upper limit, but it is preferable to keep it at 98% or less. There is no particular upper limit, and bainite may be 100% in area fraction. Furthermore, since fatigue cracks originate from the inner surface of the steel pipe, uniformity of the internal structure of the steel pipe is important. Therefore, in the case of steel pipes, the metallic structure at a position 1/4 of the wall thickness from the inner surface of the steel pipe is specified.
Here, the bainite structure includes bainitic ferrite or granular bainite that transforms during or after controlled cooling, contributing to transformational strengthening, and also includes tempered bainite. If dissimilar structures such as ferrite, martensite, pearlite, island martensite, or retained austenite are mixed in the bainite structure, a decrease in strength and deterioration of toughness will occur. Therefore, the smaller the volume fraction of structures other than the bainite phase, the better. Here, the martensite structure described above includes tempered martensite.

The hydrogen diffusion coefficient at room temperature is 1.5 × ^{10⁻¹⁰ m²} /s or higher. The fatigue crack propagation rate in hydrogen accelerates as the amount of hydrogen accumulated at the crack tip increases. The smaller the hydrogen diffusion coefficient, the greater the amount of hydrogen accumulated at the crack tip, and the greater the fatigue crack propagation rate. If the hydrogen diffusion coefficient is smaller than 1.5 × ^{10⁻¹⁰ m²} /s, the fatigue crack propagation rate in hydrogen increases significantly, so the hydrogen diffusion coefficient was set to 1.5 × ^{10⁻¹⁰ m²} /s or higher. Preferably, the hydrogen diffusion coefficient is 2.0 × ^{10⁻¹⁰ m²} /s or higher, and more preferably 3.0 × ^{10⁻¹⁰ m²} /s or higher. Even more preferably, the hydrogen diffusion coefficient is 5.0 × ^{10⁻¹⁰ m²} /s or higher. Most preferably, the hydrogen diffusion coefficient is 6.0 × ^{10⁻¹⁰ m²} /s or higher. While there is no particular upper limit, reducing the hydrogen diffusion coefficient comes with a decrease in strength, so considering the material strength, a value of 5.0 × ^{10⁻⁹ m²} /s or less is preferable. Since retained austenite has a low hydrogen diffusion coefficient at room temperature, the retained austenite fraction described above is necessary to achieve the above value for the hydrogen diffusion coefficient at room temperature. Furthermore, hydrogen penetration into steel materials occurs from the surface of the steel material (or the inner surface of the steel pipe in the case of a steel pipe). Therefore, the important value of this hydrogen diffusion coefficient is the value up to the thickness at which fatigue cracks propagate through the wall thickness and lead to rapid fracture. The thickness at which rapid fracture occurs can be determined by the fracture toughness value of the material and the stress generated in the steel pipe. However, in reality, the fatigue crack propagation life of steel materials (or steel pipes in the case of steel pipes) is mostly taken up to the point where the crack propagates through 1/4 t of the wall thickness, so the hydrogen diffusion coefficient only needs to be measured from the inner surface of the steel pipe at 1/4 t. Furthermore, since the hydrogen diffusion coefficient is highly temperature-dependent, the hydrogen diffusion coefficient of this invention is defined as the value at room temperature (20 ± 10°C).

Hydrogen solid solubility is 0.05 mass ppm/√P or less. In this invention, hydrogen solid solubility is the most important factor. The fatigue crack propagation rate is greatly influenced by the accumulation of hydrogen at the crack tip and the stress (stress intensity factor) at the crack tip. In other words, in order to obtain a desired fatigue crack propagation rate, it is important to reduce the accumulation of hydrogen at the crack tip. The smaller the hydrogen solid solubility, the greater the decrease in the fatigue crack propagation rate in hydrogen, and in order to obtain a desired fatigue crack propagation rate, the hydrogen solid solubility in the steel material should be 0.05 mass ppm/√P or less. 0.03 mass ppm/√P or less is preferable, and 0.02 mass ppm/√P or less is more preferable. Stabilization treatment or dehydrogenation treatment is carried out to reduce the hydrogen solid solubility. On the other hand, heat treatment that reduces the hydrogen solid solubility to less than 0.005 mass ppm/√P leads to a decrease in material strength and a significant increase in manufacturing costs; therefore, a hydrogen solid solubility of 0.005 mass ppm/√P or higher is preferable.

Here, the hydrogen solid solubility s is a value that represents the slope [mass ppm/√P] between the amount of hydrogen H [mass ppm] entering in an environment with hydrogen pressure P [MPa] and the square root of the hydrogen pressure P [MPa]. In the case of a mixed gas environment, P can be read as the hydrogen partial pressure equivalent P'. Specifically, in the case of a gas environment containing 20% hydrogen in a 25 MPa environment, the hydrogen partial pressure P' is 25 MPa × 0.2 = 5 MPa.

There are several methods for calculating hydrogen solid solubility, and one example is shown below. For instance, a test specimen is exposed to an arbitrary pressure environment of 0 to 40 MPa high-pressure hydrogen gas and held for a predetermined time. Then, the amount of hydrogen in the steel is measured using a hydrogen analyzer, the relationship between H and √P is obtained, and s is calculated from its slope. Alternatively, it can be calculated by a cathode hydrogen charge test that simulates a high-pressure gas environment, such as in Non-Patent Document 2.

Furthermore, the tensile strength is preferably 520 MPa or higher, and more preferably 580 MPa or higher. While there is no particular upper limit, the tensile strength is preferably 950 MPa or lower, and more preferably 800 MPa or lower.

Furthermore, although not particularly limited, a plate thickness of 5 mm or more is preferred. A plate thickness of 30 mm or less is preferred.

Next, the method for manufacturing the steel sheet of the present invention will be described. The steel material of the present invention can be manufactured by sequentially performing a heating process of the steel material (slab), a hot rolling process, a controlled cooling process, and then either a stabilization process or a dehydrogenation process.

In the following explanation, unless otherwise specified, the temperature refers to the temperature at the center of the thickness of the steel material or steel pipe. The average cooling rate refers to the temperature at a point 1/4 of the way through the wall thickness from the inner surface of the steel pipe. The temperature at the center of the thickness and the temperature at a point 1/4 of the way through the wall thickness from the inner surface of the steel pipe are estimated using heat transfer calculations that take into account the heat transfer coefficient of the steel material, based on the surface temperature of the steel pipe measured with a radiation thermometer.

Heating temperature for steel material: 1000-1250°C
When heating steel materials such as billets and slabs to temperatures below 1000°C, the diffusion of micro-segregated impurity elements such as C, P, and S is insufficient, resulting in an unholy material. For this reason, the heating temperature of the steel material should be 1000°C or higher. Preferably, the heating temperature of the steel material should be 1180°C or higher, and more preferably 1200°C or higher. On the other hand, if the temperature exceeds 1250°C, the crystal grains become too coarse, and the toughness deteriorates. Therefore, the heating temperature of the steel material should be 1250°C or lower. Preferably, the heating temperature of the steel material should be 1230°C or lower, and more preferably 1210°C or lower.

Hot Rolling Termination Temperature: Ar ₃ or higher After reheating the steel material, it is hot-rolled to the desired pipe or plate thickness, but the termination temperature of the hot rolling should be Ar ₃ or higher, which is the ferrite formation temperature. If the temperature is below Ar ₃ , in the case of a process in which cooling is performed immediately after hot rolling, the formation of a soft ferrite phase will lead to a decrease in strength. The termination temperature of the hot rolling should preferably be 770°C or higher, and if Ar ₃ is higher than 770°C, the termination temperature of the finish rolling should preferably be Ar ₃ + 30°C or higher, and more preferably Ar ₃ + 50°C or higher. Furthermore, if the temperature exceeds 1250°C, the grains become too coarse and the toughness deteriorates, so it is preferable to keep the upper limit at 1250°C or lower.

Since the Ar _3- point temperature varies depending on the alloy composition of the steel, it can be determined by experimentally measuring the transformation temperature of each steel, but it can also be determined from the component composition using the following formula.
Ar ₃ (℃) = 910-310C (%) - 80Mn (%) - 20Cu (%) - 15Cr (%) - 55Ni (%) - 80Mo (%)
Each alloying element is given as its content (mass %).

Controlled Cooling Process Cooling Start Temperature: Steel plate surface temperature of Ar ₃ or higher If the steel plate surface temperature at the start of cooling is less than Ar ₃ , ferrite will form before controlled cooling, resulting in a significant decrease in strength. For this reason, the steel plate surface temperature at the start of cooling should be Ar ₃ or higher. Preferably, the steel plate surface temperature at the start of cooling should be 770°C or higher. If Ar ₃ is higher than 770°C, the finish rolling end temperature should preferably be Ar ₃ + 30°C or higher, and more preferably Ar ₃ + 50°C or higher. There is no particular upper limit, but it is preferable to keep it at 1250°C or lower. Note that the steel plate surface temperature at the start of cooling is the temperature at the tail end of the steel plate where the cooling start temperature is lowest.

Time difference between the leading and trailing ends of the steel sheet: within 50 seconds. If the time difference between the leading and trailing ends in the rolling direction of the steel sheet at the start of cooling exceeds 50 seconds, the temperature difference between the leading and trailing ends at the start of cooling will be large, resulting in large temperature variations when cooling stops, large variations in Vickers hardness at 0.25 mm below the surface of the steel sheet, and deterioration of HISC resistance. For this reason, the time difference between the leading and trailing ends of the steel sheet at the start of cooling should be within 50 seconds, preferably within 45 seconds. More preferably within 40 seconds. It is possible to shorten the time difference at the start of cooling by shortening the length of the steel sheet, but this reduces manufacturability, so it is preferable to shorten the time difference at the start of cooling by increasing the steel sheet conveying speed. There is no particular lower limit, but the time difference at the start of cooling may be greater than 0 seconds.

Cooling Rate in Controlled Cooling Process: In order to achieve high strength while obtaining excellent resistance to HISC (High-Intensity Stabilization), it is necessary to control the cooling rate at 0.25 mm below the surface of the steel plate and at the center of the plate thickness. The cooling rate in the thickness direction is a value obtained by simulation using heat transfer calculations etc. from the surface temperature measured with a radiation thermometer.

Average cooling rate from 750°C to 550°C at a depth of 0.25 mm below the surface of the steel plate: 15–50°C/s
It is important to minimize the average cooling rate of the steel plate temperature at a depth of 0.25 mm below the surface, from 750°C to 550°C, in order to create granular bainite. Since the temperature range from 750°C to 550°C is a crucial temperature range for bainite transformation, it is important to control the cooling rate in this temperature range. If the cooling rate exceeds 50°C/s, variations in hardness may occur, and the HISC resistance after pipe manufacturing will deteriorate. Therefore, the average cooling rate should be 50°C/s or less. Preferably, it should be 45°C/s or less. More preferably, it should be 40°C/s or less. On the other hand, if the cooling rate is too low, ferrite and pearlite will be formed, resulting in insufficient strength. From the viewpoint of preventing this, it is preferable to have a cooling rate of 15°C/s or more, and preferably 17°C/s or more. More preferably, it should be 20°C/s or more, and even more preferably 25°C/s or more. Furthermore, for cooling to a steel plate temperature of 550°C or lower at a depth of 0.25 mm below the surface of the steel plate, if the cooling rate is slow, cooling may not occur in a stable nucleated boiling state, and there is a risk of hardness variations in the outermost layer of the steel plate. Therefore, it is preferable that the average cooling rate from a steel plate temperature of 550°C at a depth of 0.25 mm below the surface of the steel plate to the cooling stop temperature be 150°C/s or higher. To avoid the risk of hardness variations, it is preferable that the average cooling rate be 250°C/s or lower.

Average cooling rate from 750°C to 550°C at the center of the plate thickness: 15–50°C/s
If the average cooling rate from 750°C to 550°C in the center of the plate thickness is less than 15°C/s, a granular bainite structure cannot be obtained, resulting in a decrease in strength. In addition, excessive retained austenite is produced, and the hydrogen diffusion coefficient at room temperature decreases. For this reason, the average cooling rate in the center of the plate thickness should be 15°C/s or higher. From the viewpoint of suppressing variations in structure, it is preferable that the average cooling rate in the center of the plate thickness be 17°C/s or higher. More preferably, the average cooling rate in the center of the plate thickness is 20°C/s or higher, and even more preferably 25°C/s or higher. On the other hand, in order to suppress variations in grain size, it is preferable that the average cooling rate be 50°C/s or lower, and more preferably 45°C/s or lower. More preferably, the average cooling rate in the center of the plate thickness is 40°C/s or lower. Note that there are no particular limitations on cooling to a steel plate temperature of 550°C or lower in the center of the plate thickness, but from the viewpoint of suppressing variations in structure and grain size, it is preferable that the average cooling rate in the center of the plate thickness be 15°C/s or higher. Furthermore, it is preferable that the average cooling rate at the center of the plate thickness be 50°C/s or less.
Furthermore, if the carbon content is high, the transformation morphology changes from bainite transformation to martensitic transformation. However, if the average cooling rate from 750°C to 550°C in the center of the cooled plate thickness is less than 15°C/s, a mixed structure of martensite and bainite will be formed. Therefore, the average cooling rate should be 15°C/s or higher. From the viewpoint of suppressing variations in structure, it is preferable that the average cooling rate in the center of the plate thickness be 17°C/s or higher. More preferably, the average cooling rate in the center of the plate thickness is 20°C/s or higher, and even more preferably 25°C/s or higher. On the other hand, in order to suppress variations in grain size, it is preferable that the average cooling rate be 50°C/s or lower, and more preferably 45°C/s or lower. More preferably, the average cooling rate in the center of the plate thickness is 40°C/s or lower. Note that there are no particular limitations on cooling to a steel plate temperature of 550°C or lower in the center of the plate thickness, but from the viewpoint of suppressing variations in structure and grain size, it is preferable that the average cooling rate in the center of the plate thickness be 15°C/s or higher. Furthermore, it is preferable that the average cooling rate in the center of the plate thickness be 50°C/s or lower.

Although the steel plate temperature at 0.25 mm below the surface and at the center of the plate thickness cannot be directly measured physically, it can be calculated in real time from the temperature distribution within the plate thickness cross-section by differential calculation using a process computer, for example, based on the surface temperature at the start of cooling and the surface temperature at the target cooling stop, as measured by a radiation thermometer. The temperature at 0.25 mm below the surface of the steel plate in this temperature distribution is referred to as the "steel plate temperature at 0.25 mm below the surface of the steel plate" in this specification, and the temperature at the center of the plate thickness in this temperature distribution is referred to as the "steel plate temperature at the center of the plate thickness" in this specification.

Cooling stop temperature: The temperature of the steel plate at 0.25 mm below the surface and at the center of the plate thickness is 250 to 650°C.
If the cooling stop temperature exceeds 650°C, the bainite transformation becomes incomplete, and sufficient strength cannot be obtained. For this reason, the cooling stop temperature should be 650°C or lower. Preferably, the cooling stop temperature should be 625°C or lower, more preferably 600°C or lower, and even more preferably 500°C or lower. Also, if the cooling stop temperature is below 250°C, the hardness increases, and the resistance to HISC deteriorates. For this reason, the cooling stop temperature should be 250°C or higher. Preferably, the cooling stop temperature should be 270°C or higher. More preferably, the cooling stop temperature should be 300°C or higher.

Tempering Process Tempering treatment may be performed for the purpose of improving toughness or adjusting material strength. Since the effect of tempering cannot be obtained below 200°C, it is preferable to set the tempering temperature to 200°C or higher if performed. On the other hand, tempering can also cause a decrease in strength, and if the temperature becomes too high, the microstructure will undergo another transformation, so it is preferable to keep the Ar point below _3. The holding time can be arbitrarily determined, but it is preferable to hold it for 10 minutes or more at a predetermined temperature in the center of the plate thickness. It is preferable to hold it for 180 minutes or less.

Stabilization Process: Hydrogen that penetrates steel is mainly trapped in various defects such as dislocations. When hydrogen is trapped in these defects, the hydrogen diffusion coefficient decreases and the hydrogen solid solubility increases. As a result, the hydrogen embrittlement resistance deteriorates. Therefore, it is important to reduce these defects or reduce the binding of hydrogen to these defects. For this reason, a dislocation stabilization treatment is performed after manufacturing to weaken the binding of hydrogen to dislocations. By holding the product at a predetermined temperature for a certain period of time before use, solid-solution carbon can be fixed to the dislocations, and by stabilizing the dislocations, the binding of hydrogen to dislocations can be reduced. As a result, it is possible to increase the hydrogen diffusion coefficient and decrease the hydrogen solid solubility, making it possible to obtain steel with excellent hydrogen embrittlement resistance in a high-pressure hydrogen gas environment.
The stabilization process is performed before pipe manufacturing and welding to connect steel pipes. The temperature should be above room temperature (25°C ± 10°C) because carbon diffusion is significantly low below this temperature. Furthermore, a temperature of 100°C or higher is preferable, and 200°C or higher is more preferable, as higher temperatures result in a smaller carbon diffusion coefficient Dc and carbon diffuses in a shorter time. On the other hand, if the temperature of the stabilization process is too high, the material strength will decrease significantly, so the stabilization temperature should be below the Ar ₃ point (°C) or below 700°C. Also, when performing stabilization on tempered material, it is preferable to set the upper limit to a temperature at least 50°C lower than the tempering temperature. The holding time should be 72 hours or more if the stabilization temperature is below 100°C, and 10 minutes or more if it is 100°C or higher. The holding time should preferably be 400 hours or less if the stabilization temperature is below 100°C, and 100 hours or less if it is 100°C or higher. The temperature should be centered on the plate thickness.
Furthermore, the time and temperature of the stabilization process may be combined with the heating process in the pipe manufacturing process for electric resistance welded pipes, UOE steel pipes, etc., if such heating is performed during the pipe manufacturing process. This process refers to processes that perform heat treatment after pipe manufacturing, such as tempering or stress-relieving annealing.

Dehydrogenation Treatment Process: When hydrogen is present in steel, the acceleration of fatigue crack propagation increases, reducing fatigue life and hydrogen fatigue limit stress. Therefore, dehydrogenation treatment may be used to release any remaining hydrogen after manufacturing. Dehydrogenation treatment reduces the amount of hydrogen in the steel by holding it at a high temperature for a certain period of time before product use, making it possible to obtain steel with excellent resistance to hydrogen embrittlement in a high-pressure hydrogen gas environment.
The holding time R (sec) is preferably given by the following formula (A), based on the plate thickness and pipe thickness t (mm) of the steel material and steel pipe, and the hydrogen diffusion coefficient D ( ^mm² · ^sec⁻¹ ) in steel at room temperature.
R≧t ² /D...(A)
Furthermore, the hydrogen diffusion coefficient can be calculated using the method described above.
The dehydrogenation treatment process is carried out before pipe fabrication and welding to connect steel pipes. It is preferable to perform the dehydrogenation treatment at a high temperature because the hydrogen diffusion coefficient D decreases at high temperatures, allowing hydrogen to escape more quickly. At high temperatures, the diffusion coefficient D' (diffusion coefficient at each temperature) at which the value of D in equation (A) above is maintained may be used for calculation. Furthermore, if the temperature T of the dehydrogenation process is too high, the material strength will decrease significantly, so the dehydrogenation treatment temperature is preferably 550°C or lower. It is more preferable that the dehydrogenation treatment temperature T be 500°C or lower. It is even more preferable that the dehydrogenation treatment temperature T be 400°C or lower, and most preferably 300°C or lower. Also, it is preferable that the dehydrogenation treatment temperature T be above room temperature because dehydrogenation treatment at temperatures lower than room temperature increases processing time and cost. It is more preferable that the dehydrogenation treatment temperature T be 50°C or higher. It is even more preferable that the dehydrogenation treatment temperature T be 100°C or higher, and most preferably 150°C or higher. The dehydrogenation treatment temperature T is the temperature of the atmosphere during the dehydrogenation treatment process. Room temperature is defined as 20 ± 10°C.
Furthermore, when performing dehydrogenation treatment on tempered materials, the upper limit should be at least 50°C lower than the tempering temperature.

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 ambient temperature (dehydrogenation treatment temperature T) in the dehydrogenation treatment process. Therefore, even if the above holding time R (sec) is met at the ambient temperature, if the center of the thickness has not reached the dehydrogenation treatment temperature T (ambient temperature), the dehydrogenation treatment may be insufficient. For this reason, it is preferable to hold the temperature Tc at the center of the thickness for R (sec) or longer after it reaches the target temperature T. Furthermore, in order to obtain a predetermined crack propagation rate in hydrogen gas, it is necessary to appropriately adjust the amount of hydrogen in the steel material at the surface and the center of the thickness. For this purpose, it is preferable to hold the temperature at the ambient temperature T for R (sec) or longer as defined by equation (A), and it is even more preferable to hold the temperature Tc at the center of the thickness for R (sec) or longer after it reaches the target temperature T. In other words, at least the former allows for appropriate control of the amount of hydrogen in the surface of the steel material and steel pipe, and if the latter is also implemented, the amount of hydrogen in the steel material from the surface to the center of the thickness can be appropriately controlled. The central temperature Tc of the plate thickness can be measured using thermocouples or other methods, or it can be predicted using methods such as the finite element method.

Furthermore, the time and temperature of the dehydrogenation treatment process may include the temperature and time applied during the heating process in the pipe manufacturing process, such as for electric resistance welded pipes or UOE pipes, as described later. Additionally, since scale on the steel surface inhibits dehydrogenation, it is preferable to remove the scale before performing the dehydrogenation treatment. The removal method is not limited; for example, physical cleaning by high-pressure washing or a chemical method using a scale remover may be used. A removal of approximately 100 μm in thickness is sufficient to achieve the desired scale removal effect.

Second Embodiment Furthermore, UOE steel pipes, which can be cited as an example of high-strength steel pipes for line pipes, can be obtained by limiting the manufacturing conditions shown below, and the manufacturing method and conditions will be explained in detail. The component composition, metal structure, hydrogen solid solubility, and hydrogen diffusion coefficient of UOE steel pipes are the same as those described for the steel material in the first embodiment, and the manufacturing method, including the heating process, hot rolling process, controlled cooling process after hot rolling, stabilization process, and dehydrogenation process, is carried out in the same manner as described for the steel material. The pipe-making process after rolling will be explained in detail below.

UOE steel pipes are manufactured by bending hot-rolled steel sheets, specifically by beveling the ends of the hot-rolled steel sheets, forming them into a steel pipe shape using C-press, U-press, and O-press, then seam welding the butt joints using internal and external welding, and further expanding the pipe as needed. Any welding method is acceptable as long as sufficient joint strength and toughness can be obtained, but submerged arc welding is preferred from the viewpoint of excellent welding quality and manufacturing efficiency. Furthermore, pipe expansion can also be performed on steel pipes that have been formed into a tubular shape by press bending and then seam-welded at the butt joints.

Third Embodiment Furthermore, as an example of the high-strength steel pipe for line pipes according to the present invention, electric resistance welded (ERW) steel pipes can be obtained by limiting the manufacturing conditions shown below, and the manufacturing method and conditions will be explained in detail. The component composition, metal structure, hydrogen solid solubility, and hydrogen diffusion coefficient of the steel material are the same as those described for the steel material in the first embodiment, and the manufacturing method is also carried out in the same manner as described for the steel material, except for the controlled cooling process after rolling and the pipe making process (heating process, hot rolling process, stabilization process, dehydrogenation process).

Cooling process after rolling (controlled cooling process)
The cooling start temperature and average cooling rate of the controlled cooling are the same as those described in the first embodiment.

Cooling stop temperature: 250-650℃
If the cooling stop temperature after hot rolling exceeds 650°C, the bainite transformation becomes incomplete, and the material strength decreases significantly. For this reason, the cooling stop temperature should be 650°C or lower. Preferably, the cooling stop temperature should be 620°C or lower. More preferably, the cooling stop temperature should be 600°C or lower, and even more preferably 580°C or lower. On the other hand, if the cooling stop temperature is below 250°C, quench cracking during cooling is likely to occur. Also, in order to obtain a uniform bainite structure, the cooling stop temperature should be 250°C or higher. From the standpoint of suppressing the amount of hydrogen in the steel, it is also necessary to set the cooling stop temperature above a predetermined temperature. Specifically, hydrogen present in the steel gradually escapes during cooling, and the effect is greater at higher temperatures, but if the cooling stop temperature is too low, supercooling occurs, and hydrogen remains in the steel. Furthermore, if the cooling stop temperature is set too low, retained austenite, which rapidly increases in hydrogen compared to other phases, is likely to form. For this reason, the cooling stop temperature should be 250°C or higher in order to reduce the amount of hydrogen in the steel. The cooling stop temperature is preferably 300°C or higher, more preferably 390°C or higher. Even more preferably, the cooling stop temperature is 450°C or higher. After cooling stops, it is fine to let it cool naturally, but to promote the formation of bainite, it is more preferable to cool it slowly until the temperature drops by about 50°C from the cooling stop temperature. Note that the cooling stop temperature referred to here is the temperature at the center of the plate thickness.

Subsequently, the hot-rolled steel sheet obtained as described above is wound into a coil. The winding temperature is preferably 550°C or lower.

Pipe Manufacturing Process An electric resistance welded (ERW) steel pipe, as an example of the present invention, is manufactured by forming it into a cylindrical shape by cold roll forming and welding the circumferential ends of the cylindrical shape together. Furthermore, it may also be manufactured by forming it into an ERW steel pipe material (ERW steel pipe) using a sizing roll that satisfies the following equation (1) (sizing process), and then applying an internal pressure p (MPa) that satisfies the following equation (2) to the inner surface of the ERW steel pipe material (internal pressure loading process).
Furthermore, the term "cylindrical" refers to a pipe whose circumferential cross-section is "C" shaped.
Diameter of sizing roll (mm) ≥ Thickness of hot-rolled steel sheet (mm) / 0.020 ... (1)
The thickness of a hot-rolled steel sheet refers to the thickness of the hot-rolled steel sheet before the sizing process.
X<p≦X×1.5 (2)
Note that X = (wall thickness of electric resistance welded steel pipe material (mm) / radius of electric resistance welded steel pipe material (mm)) × yield strength of electric resistance welded steel pipe material (MPa)
The aforementioned internal pressure load can be implemented, for example, by sealing the pipe end with a rubber gasket and applying water pressure inside the pipe. Furthermore, to stabilize the shape, a mold of the desired diameter can be used as an outer frame if necessary.

Furthermore, the wall thickness of the electric resistance welded (ERW) steel pipe material, which is given as an example of the steel pipe of the present invention, is preferably 5 mm or more, and preferably 30 mm or less. While there is no upper limit specified for the radius of the ERW steel pipe material, a larger radius increases the load on the equipment, so a radius of 400 mm or less is preferred. A radius of 200 mm or more is also preferred. In addition, the yield strength of the ERW steel pipe material is preferably 480 MPa or more to withstand the gas pressure of pipeline operation. A yield strength of 500 MPa or more is more preferred. On the other hand, to avoid increased susceptibility to hydrogen embrittlement, a yield strength of 600 MPa or less is preferred. A yield strength of 560 MPa or less is more preferred.

During the sizing process, bending deformation occurs in the axial direction of the pipe along the roll shape as the pipe passes through the roll, generating residual stress in the axial direction of the pipe. The greater the bending strain in the bending deformation, the greater the absolute value of the residual stress in the axial direction of the pipe. The bending strain increases as the diameter of the sizing roll decreases and as the thickness of the hot-rolled steel sheet increases.
Therefore, in this invention, from the viewpoint of reducing shear residual stress, the diameter of the sizing roll is set to satisfy equation (1) above in order to reduce the absolute value of residual stress in the axial direction of the pipe.
If the diameter of the sizing roll is less than the right-hand side of equation (1) above, the shear residual stress targeted by the present invention cannot be obtained. Although there is no upper limit specified for the diameter of the sizing roll, it is preferable that the diameter of the sizing roll be 2000 mm or less, as a larger sizing roll increases the load on the equipment.

In the internal pressure loading 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 residual stress in the circumferential direction.
The greater the internal pressure p (MPa) during the internal pressure loading process, the smaller the absolute value of the residual stress in the circumferential direction of the pipe. The tensile stress generated in the circumferential direction of the pipe increases as the radius of the steel pipe increases and as the wall thickness of the steel pipe decreases.

The left-hand side (X) of equation (2) above 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 this invention, from the viewpoint of reducing shear residual stress, the internal pressure p is set to a value greater than (X) on the left side of equation (2) to reduce the absolute value of residual stress in the axial direction of the pipe, and the electric resistance welded steel pipe material is expanded to the plastic region. On the other hand, if the internal pressure p exceeds (X × 1.5) on the right side of equation (2), the absolute value of residual stress in the circumferential direction of the pipe decreases, but the amount of work hardening due to pipe expansion becomes too large, the dislocation density on the pipe surface increases, and the hydrogen embrittlement resistance decreases.

As partially explained above, the steel pipes of the present invention can be manufactured by forming the steel material disclosed in this invention into a tubular shape using press bending, roll forming, UOE forming, etc., and then welding the butt joints. This allows for the production of high-strength steel pipes for sour line pipes (UOE steel pipes, electric resistance welded steel pipes, spiral steel pipes, etc.) with excellent material uniformity within the steel plate, suitable for the transportation of crude oil and natural gas. Furthermore, by using the steel material disclosed in this disclosure for steel pipes, it is possible to manufacture steel pipes with excellent HISC resistance even if a high-hardness region exists in the welded area.

Next, the present invention will be described in more detail based on examples. The following examples illustrate preferred examples of the present invention, and the present invention is not limited in any way by the examples described.

Steel pipes were manufactured from steel materials with the component composition shown in Table 1. The manufacturing procedure was as follows: First, billets with the component composition shown in Table 1 were prepared. The casting speed was 0.05 to 0.2 m/min. The billets were heated to 1000°C to 1100°C and hot-rolled in the range of 950 ± 50°C. Controlled cooling was started when the surface temperature reached 900°C. The difference in cooling start time between the leading and trailing ends of the hot-rolled pipe was 30 to 45 seconds, and the cooling stop temperature was set to 300 ± 50°C. The target thickness of the steel plate was 20 mm. The average cooling rate during the controlled cooling process was carried out under the conditions shown in Table 2. For some steel materials (steel materials No. 1 to 11), after the controlled cooling process, a pipe-making process was performed in which the hot-rolled steel sheets were bent and both ends were butt-welded. For some steel materials (steel materials No. 12 to 22), after the controlled cooling process, the hot-rolled steel sheets were formed into a cylindrical shape by cold roll forming, and the circumferential ends of the cylindrical shape were butt-welded using electric resistance welding. In this way, steel pipes No. 1 to 22 were obtained.
Subsequently, a circle (○) was marked for materials that underwent stabilization treatment (or dehydrogenation treatment). All treatment conditions were 200°C for 30 minutes. The steel materials and pipes manufactured as described above were evaluated according to the following criteria.

Furthermore, billets with the component compositions shown for steel grade No. 8 and steel grade No. 22 in Table 1 were produced at various casting speeds shown in Table 3. The billets were heated to 1000°C to 1100°C, and hot rolling was performed in the range of 950±50°C. Controlled cooling was started when the surface temperature reached 900°C. The difference in cooling start time between the leading and trailing ends of the hot rolling was 30 to 45 seconds, and the cooling stop temperature was set to 300±50°C. The target thickness of the steel plate was 20 mm. The average cooling rate in the controlled cooling process was carried out under the conditions shown in Table 3, and steel materials and steel pipes were obtained. Steel materials No. 8-1, 8-2, 8-3, 22-1, 22-2, and 22-3 were in their raw state, while steel pipe No. 8-11, 8-12, and 8-13 were manufactured by bending hot-rolled steel sheets and welding the ends together in a pipe-making process. Steel pipes No. 22-11, 22-12, and 22-13 were obtained by cold-rolling hot-rolled steel sheets into a cylindrical shape after a controlled cooling process, and then welding the ends of the cylindrical shape together in a pipe-making process using electric resistance welding. Subsequently, a stabilization treatment (or dehydrogenation treatment) was performed, indicated by a circle (〇), and the treatment conditions for all of these were 200°C for 30 minutes. The steel materials and steel pipes manufactured as described above were evaluated according to the following criteria.

Measurement of Area Fraction of Retained Austenite Samples for metallographic observation were taken from the center of the plate width in the longitudinal direction of the steel material and steel pipe obtained according to the above procedure. The cross section parallel to the longitudinal direction was used as the observation surface and polished with a buff. After that, the surface layer was removed by chemical polishing using picric acid etching, and the area fraction was measured using X-ray diffraction. Specifically, a Co-Kα source was used for the incident X-rays, and the area fraction of retained austenite was calculated from the intensity ratio of the (200), (211), (220) planes of ferrite and the (200), (220), (311) planes of austenite.

Area fraction measurement of bainite and martensite The metallographic structure at the 1/4 wall thickness position on the inside of the obtained steel pipe was evaluated as follows: Test specimens were taken from the longitudinal center of the steel pipe so that the 1/4 wall thickness position and the wall thickness center position on the inside were the observation positions, and the cross-sections of the taken test specimens were etched with a 3 vol% nital solution. Scanning electron microscope (STEM) images were taken at an appropriate magnification between 1000 and 5000x to observe martensite (including tempered martensite), ferrite, bainite, and pearlite. Martensite and bainite were identified visually by comparing them with the tissue photographs in Non-Patent Document 3. The tissue fractions were determined by image analysis using images obtained by dividing SEM photographs based on the above identification (for example, to calculate the fraction of bainite, the bainite region and the other regions were binarized to determine the bainite fraction), and these were taken as the area fractions of each phase.

Hydrogen Temperature Analysis The amount of hydrogen remaining in the steel was determined using a temperature-controlled desorption analysis method with a low-temperature temperature-controlled hydrogen analyzer (gas chromatograph type) (JTF-20AL). The temperature-controlled desorption analysis was performed at a heating rate of 200°C/h in the temperature range from room temperature to 400°C, and the sum was taken as the hydrogen amount. The test specimens were rectangular prisms with a length of 20 mm in the longitudinal direction of the steel pipe, 10 mm thick and 10 mm wide, located at the 1/4 position of the thickness of the steel plate and at the 1/4 position from the inner surface of the steel pipe. Note that this hydrogen amount was obtained before subjecting the specimens to the high-pressure hydrogen fatigue crack propagation test and high-pressure hydrogen exposure test described in the next section.

High-pressure hydrogen exposure test (calculation of hydrogen solid solubility)
This section describes the method for calculating hydrogen solid solubility. First, the test specimens were prismatic in shape, measuring 20 mm in length and 10 mm in thickness × 10 mm in width, positioned at 1/4 of the thickness of the steel plate and 1/4 of the inner surface of the steel pipe, along the longitudinal direction of the steel and steel pipe. Since the amount of hydrogen penetration depends on the surface condition of the test specimen, after cutting, the specimens were polished with emery paper from 160 to 1000 grit to ensure a uniform surface condition for all samples. Subsequently, Pd plating was applied to the entire surface of the test specimens to remove oxide films that may inhibit hydrogen penetration. This Pd plating can be done by other methods such as vapor deposition, and Ni plating can also be used as a substitute.
The above test specimens were exposed to a high-pressure hydrogen environment (over 99.999% hydrogen by volume fraction) at room temperature (20 ± 10°C) and pressures of 0, 5, 25, and 40 MPa for 72 hours. After exposure, the specimens were promptly removed from the exposure environment and stored in liquid nitrogen to prevent hydrogen from being released from the specimens. The amount of absorbed hydrogen, H, was determined using the same hydrogen temperature analysis method as described above. The x-axis plots the square root of the exposure pressure, √P [MPa], and the y-axis plots the measured hydrogen absorbed amount, H [mass ppm/√P]. The amount of hydrogen in the specimen at 0 MPa (before exposure testing) was used as the initial hydrogen amount (intercept), and the hydrogen solid solubility s [mass ppm/√P] was calculated from the slope of √P - H.

Hydrogen Diffusion Coefficient The hydrogen diffusion coefficient was evaluated using 1 × 40 × 40 mm test specimens taken from the center of the steel plate thickness at a position 1/4 of the plate thickness and at a position 1/4 of the inner surface of the steel pipe. One side of the test specimen was plated with Ni, and using a Devanathan-type cell, the unplated side was immersed in a 0.2% NaCl solution for cathode hydrogen charging, and the Ni-plated side was immersed in a 0.1N NaOH aqueous solution to set the withdrawal potential to 0V. The hydrogen permeation start time (2nd Build up), which is the rise of the second permeation current, was fitted to the theoretical curve in Non-Patent Document 4 to determine the diffusion coefficient.

Evaluation of High-Pressure Hydrogen Fatigue Crack Propagation Rate The fatigue crack propagation rate was determined by conducting fatigue tests in a mixed atmosphere of hydrogen gas at room temperature (20 ± 10°C), pressure: 25 MPa, hydrogen gas at a pressure of 1 MPa or higher, or natural gas containing hydrogen as a partial pressure of 1 MPa or higher (main components being hydrocarbons such as methane and ethane), in accordance with ASTM E647, Fatigue Testing, with a frequency of 1 Hz, repetition waveform: sine wave, control method: load control, load condition: uniaxial tension, and stress ratio: R = 0.1. Steel materials and steel pipes with excellent hydrogen embrittlement resistance are defined as those with a hydrogen fatigue crack propagation rate da/dNmm/cycle of 2.0 × ^10⁻³ mm/cycle or less at ΔK = 25 MPa.

Tensile strength (TS)
From the steel materials and steel pipes obtained according to the above, JIS No. 14 proportional test specimens (parallel section diameter 7 mm, gauge length 35 mm) were taken in accordance with JIS Z 2201, and their tensile strength was measured.

Steel materials and steel pipes satisfying the present invention example showed excellent effects on fatigue crack propagation characteristics in a hydrogen environment. Furthermore, when the hydrogen solid solubility s was less than 0.02 mass ppm/√P, the fatigue crack propagation rate was reduced to 1.5 × ^10⁻³ mm/cycle or less, an improvement of more than 30% compared to materials with a hydrogen solid solubility s of approximately 0.05 mass ppm/√P, demonstrating excellent effects.

The following describes examples that verify the effects of the present invention. In the following examples, steel pipes were manufactured under the following manufacturing conditions and their characteristics were evaluated. Using steel grades No. 1, 8, 10, 12, and 22 shown in Tables 1-1 and 1-2, the pipes were manufactured under the same conditions as steel grades No. 1, 8-2, 10, 12, and 22-2 shown in Tables 2 and 3 up to the controlled cooling process, and their characteristics were evaluated when the dehydrogenation treatment conditions were changed. Steel pipe forming was carried out in the same manner as in Example 1. The results are shown in Table 4.

In this embodiment, the dehydrogenation treatment temperature T (ambient temperature) for steel pipes and steel materials No. 1A, 10A, 12A, 8-2A, and 22-2A was set to 50°C, and the holding time tc after the plate thickness center temperature Tc reached 50°C was carried out so as to satisfy equation (A). For steel pipes and steel materials No. 10B, 12B, 8-2B, and 22-2B, the dehydrogenation treatment temperature T (ambient temperature) was set to 50°C, and the holding time tc at the dehydrogenation treatment temperature T was set to satisfy the above-mentioned equation (A). However, the holding time tc after the plate thickness center temperature Tc reached 50°C did not satisfy the above-mentioned equation (A).

For steel pipes and steel materials No. 10C, 12C, 8-2C, and 22-2C, the dehydrogenation treatment temperature T (ambient temperature) is 50°C, but neither the holding time t at ambient temperature nor the holding time tc after the central plate thickness temperature Tc reaches 50°C satisfies equation (A) described above.

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

Various evaluations were carried out using the method described in Example 1.

All of the invention examples of this present invention satisfied excellent fatigue crack propagation rates. Among them, the fatigue crack propagation rate was superior when the dehydrogenation treatment conditions were more favorable.

Claims (10)

In mass percent,
C: 0.02-0.15%,
Si: 0.01-2.0%,
Mn: 0.5-1.8%,
P: 0.0001-0.015%,
S: 0.0002-0.0015%,
Al: 0.005-0.15%,
O: 0.01% or less,
N: 0.010% or less,
H: Includes 0.02 ppm or less,
Or, furthermore,
Nb: 0 to 0.10%,
Ca: 0-0.005%,
Ni: 0-2.0%,
Ti: 0 to 0.1%,
Cu: 0 to 1.0%,
Cr: 0-1.0%,
Mo: 0 to 0.60%,
W: 0-1.0%,
V: 0-0.10%,
Zr: 0 to 0.050%,
REM: 0-0.01%,
Mg: 0 to 0.01%,
B: 0 to 0.0020%,
Hf: 0-0.2%,
Ta: 0-0.2%,
Re: 0 to 0.005%,
Sn: 0-0.3%,
Sb: Contains one or more selected from 0 to 0.3%,
It has a chemical composition in which the remainder is Fe and unavoidable impurity elements.
A thin or thick steel plate for line pipes, having a retained austenite content of 0-3% by area fraction, a hydrogen diffusion coefficient of 1.5 × ^{10⁻¹⁰ m²} /s or higher at room temperature, and a hydrogen solid solubility of 0.05 mass ppm/√P or less. The aforementioned chemical composition is expressed in mass%, and further,
Nb: 0.001 to 0.10%,
Ca: 0.0001-0.005%,
Ni: 0.01-2.0%,
Ti: 0.005-0.1%,
Cu: 0.01 to 1.0%,
Cr: 0.01-1.0%,
Mo: 0.01 to 0.60%,
W: 0.01-1.0%,
V: 0.01-0.10%,
Zr: 0.0001 to 0.050%,
REM: 0.0001-0.01%,
Mg: 0.0001-0.01%,
B: 0.0001 to 0.0020%,
Hf: 0.0001-0.2%,
Ta: 0.0001-0.2%,
Re: 0.0001-0.005%,
Sn: 0.0001-0.3%,
The steel material for line pipes according to claim 1, comprising one or more Sb selected from 0.0001 to 0.3%. The steel material for line pipes according to claim 1 or 2, wherein the plate has bainite or martensite at the 1/4 thickness position, and the bainite or martensite accounts for 90% or more of the area fraction. A method for manufacturing steel material for line pipes according to claim 1 or 2, comprising a heating step of heating a steel material having the chemical composition at 1000 to 1250°C,
A hot rolling process is performed in which the steel material heated in the above heating process is rolled at a rolling completion temperature of Ar ₃ or higher,
A controlled cooling process is performed to cool the hot-rolled steel sheet obtained in the hot-rolling process under the following conditions: the cooling start temperature is Ar ₃ or higher at the steel sheet surface temperature, the difference in cooling start time between the leading and trailing ends 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 temperatures 0.25 mm below the steel sheet surface and in the center of the sheet thickness, and the cooling stop temperature is 250 to 650°C.
A stabilization process for stabilizing the steel sheet obtained in the controlled cooling process, or a dehydrogenation process for dehydrogenating the steel sheet obtained in the controlled cooling process,
A method for manufacturing steel materials for line pipes. A method for manufacturing steel material for line pipes according to claim 3, comprising a heating step of heating a steel material having the chemical composition at 1000 to 1250°C,
A hot rolling process is performed in which the steel material heated in the above heating process is rolled at a rolling completion temperature of Ar ₃ or higher,
A controlled cooling process is performed to cool the hot-rolled steel sheet obtained in the hot-rolling process under the following conditions: the cooling start temperature is Ar ₃ or higher at the steel sheet surface temperature, the difference in cooling start time between the leading and trailing ends 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 temperatures 0.25 mm below the steel sheet surface and in the center of the sheet thickness, and the cooling stop temperature is 250 to 650°C.
A stabilization process for stabilizing the steel sheet obtained in the controlled cooling process, or a dehydrogenation process for dehydrogenating the steel sheet obtained in the controlled cooling process,
A method for manufacturing steel materials for line pipes. In steel pipes for line pipes,
In mass percent,
C: 0.02-0.15%,
Si: 0.01-2.0%,
Mn: 0.5-1.8%,
P: 0.0001-0.015%,
S: 0.0002-0.0015%,
Al: 0.005-0.15%,
O: 0.01% or less,
N: 0.010% or less,
H: Includes 0.02 ppm or less,
Or, furthermore,
Nb: 0 to 0.10%,
Ca: 0-0.005%,
Ni: 0-2.0%,
Ti: 0 to 0.1%,
Cu: 0 to 1.0%,
Cr: 0-1.0%,
Mo: 0 to 0.60%,
W: 0-1.0%,
V: 0-0.10%,
Zr: 0 to 0.050%,
REM: 0-0.01%,
Mg: 0 to 0.01%,
B: 0 to 0.0020%,
Hf: 0-0.2%,
Ta: 0-0.2%,
Re: 0 to 0.005%,
Sn: 0-0.3%,
Sb: Contains one or more selected from 0 to 0.3%,
It has a chemical composition in which the remainder is Fe and unavoidable impurity elements.
A steel pipe for line pipes having a retained austenite content of 0-3% by area fraction, a hydrogen diffusion coefficient of 1.5 × ^{10⁻¹⁰ m²} /s or higher at room temperature, and a hydrogen solid solubility of 0.05 mass ppm/√P or less. The aforementioned chemical composition is expressed in mass%, and further,
Nb: 0.001 to 0.10%,
Ca: 0.0001-0.005%,
Ni: 0.01-2.0%,
Ti: 0.005-0.1%,
Cu: 0.01 to 1.0%,
Cr: 0.01-1.0%,
Mo: 0.01 to 0.60%,
W: 0.01-1.0%,
V: 0.01-0.10%,
Zr: 0.0001 to 0.050%,
REM: 0.0001-0.01%,
Mg: 0.0001-0.01%,
B: 0.0001 to 0.0020%,
Hf: 0.0001-0.2%,
Ta: 0.0001-0.2%,
Re: 0.0001-0.005%,
Sn: 0.0001-0.3%,
Steel pipe for line pipes according to claim 6, comprising one or more Sb selected from 0.0001 to 0.3%. A steel pipe for line pipes according to claim 6 or 7, wherein the steel pipe has bainite or martensite at a position 1/4 of the wall thickness from the inner surface of the steel pipe, and the area fraction of bainite or martensite is 90% or more. A method for manufacturing steel pipes for line pipes according to claim 6 or 7,
A heating step of heating a steel material having the aforementioned chemical composition to 1000 to 1250°C,
A hot rolling process is performed in which the steel material heated in the above heating process is rolled at a rolling completion temperature of Ar ₃ or higher,
A controlled cooling process is performed to cool the hot-rolled steel sheet obtained in the hot-rolling process under the following conditions: the cooling start temperature is Ar ₃ or higher at the steel sheet surface temperature, the difference in cooling start time between the leading and trailing ends 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 temperatures 0.25 mm below the steel sheet surface and in the center of the sheet thickness, and the cooling stop temperature is 250 to 650°C.
After the controlled cooling process, a pipe-making process is performed in which the hot-rolled steel sheet is bent and both ends are butt-welded; or a pipe-making process is performed in which the hot-rolled steel sheet is formed into a cylindrical shape by cold roll forming and both ends of the cylindrical shape are butt-welded using electric resistance welding;
Either a stabilization process for stabilizing the steel pipes obtained in the pipe manufacturing process, or a dehydrogenation process for dehydrogenating the steel pipes obtained in the pipe manufacturing process,
A method for manufacturing steel pipes for line pipes. A method for manufacturing steel pipes for line pipes according to claim 8,
A heating step of heating a steel material having the aforementioned chemical composition to 1000 to 1250°C,
A hot rolling process is performed in which the steel material heated in the above heating process is rolled at a rolling completion temperature of Ar ₃ or higher,
A controlled cooling process is performed to cool the hot-rolled steel sheet obtained in the hot-rolling process under the following conditions: the cooling start temperature is Ar ₃ or higher at the steel sheet surface temperature, the difference in cooling start time between the leading and trailing ends 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 temperatures 0.25 mm below the steel sheet surface and in the center of the sheet thickness, and the cooling stop temperature is 250 to 650°C.
After the controlled cooling process, a pipe-making process is performed in which the hot-rolled steel sheet is bent and both ends are butt-welded; or a pipe-making process is performed in which the hot-rolled steel sheet is formed into a cylindrical shape by cold roll forming and both ends of the cylindrical shape are butt-welded using electric resistance welding;
Either a stabilization process for stabilizing the steel pipes obtained in the pipe manufacturing process, or a dehydrogenation process for dehydrogenating the steel pipes obtained in the pipe manufacturing process,
A method for manufacturing steel pipes for line pipes.