JP7831570B2 - Manufacturing method for steel pipes and manufacturing method for steel materials - Google Patents

Description

This invention relates to a steel pipe with excellent fatigue properties in hydrogen, a method for manufacturing the same, a steel material, 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.

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

Furthermore, the low-alloy high-strength steel described in Patent Document 3 is a Cr-Mo-based high-strength low-alloy steel. 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, it is said to become a low-alloy high-strength steel with excellent resistance to high-pressure hydrogen environment embrittlement, exhibiting excellent elongation and reduction characteristics even in a 45 MPa hydrogen atmosphere.

Furthermore, Patent Document 4 proposes a low-alloy steel for high-pressure hydrogen gas environments. The low-alloy steel described in Patent Document 4 is said to improve the carbide morphology at grain boundaries and significantly enhance resistance to hydrogen environment embrittlement by adding V, increasing the Mo content compared to existing steels, and raising the tempering temperature to utilize V-Mo carbides.

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 for high-pressure hydrogen storage. The steel material described in Patent Document 6 has a bainite-dominant structure with an area fraction of 90% or more, in which cementite with an average particle size of 50 nm or less and an average aspect ratio of 3 or less is dispersed and precipitated within the bainite.

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

The pressure inside a line pipe fluctuates during operation and undergoes periodic shutdowns, subjecting structures like line pipes to repeated stress. Therefore, fatigue failure must be considered when designing steel structures such as line pipes. 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 strength in hydrogen gas. In other words, it is difficult to achieve both the suppression of hydrogen-induced cracking in sour environments and the high fatigue strength in hydrogen gas, which has a greater impact on service life.

In view of the problems of the prior art described above, the present invention aims to provide a steel pipe with excellent fatigue characteristics in hydrogen under high-pressure hydrogen gas conditions, a method for manufacturing the same, a steel material, and a method for manufacturing the same, which are suitable for use in steel structures used under high-pressure hydrogen gas conditions, 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).

In this context, "excellent fatigue characteristics in hydrogen under high-pressure hydrogen conditions" means that the crack propagation rate da/dN in the stress intensity factor range = 20 MPa√m, obtained by fatigue testing conducted in both environments—a hydrogen gas at room temperature (20 ± 10°C) and a pressure of 1 MPa or higher, or a mixed atmosphere of natural gas containing hydrogen as a partial pressure of 1 MPa or higher (main components being hydrocarbons such as methane and ethane)—is 1.0 × ^10⁻⁶ m· ^cycle⁻¹ or less.

Furthermore, if the crack propagation rate da/dN in a hydrogen environment is 1.0 × ^10⁻⁶ m· ^cycle⁻¹ or less, it is possible to design hydrogen-resistant structural steel within the plate thickness range that can be manufactured using the manufacturing process.

The inventors of this invention have diligently researched the conditions that various steel materials must satisfy in hydrogen gas from the above perspective, and have discovered new steel pipes and steel materials with excellent fatigue characteristics in hydrogen.

This invention was developed based on these new findings and further considerations, and its gist is as follows.
[1] In mass percent,
C: 0.10-0.45%,
Si: 0.01-2.0%,
Mn: 0.3-2.0%,
Al: 0.01-0.15%,
N: 0.0005-0.008%,
P: 0.015% or less,
S: 0.0015% or less,
O: 0.01% or less,
H: 0.0010% or less,
Cu: 0 to 2.5%,
Ni: 0 to 2.5%,
Cr: 0-2.5%,
Mo: 0-2.0%,
Nb: 0 to 0.5%,
V: 0-0.5%,
Ti: 0 to 0.5%,
W: 0-2.5%,
B: 0 to 0.005%,
Sn: 0-0.3%,
Sb: 0 to 0.3%,
Ca: 0-0.01%,
Mg: 0 to 0.01%,
REM: 0~0.005%
It contains, with the remainder being Fe and unavoidable impurities, and has a component composition.
If the residual austenite content is 3% or less,
There are 20 or fewer carbides with a diameter of 200 nm or more per ¹⁰ μm².
A steel pipe exhibiting excellent fatigue properties in hydrogen, where the crack propagation rate da/dN in hydrogen at stress intensity of 1 MPa or higher (stress intensity factor range = 20 MPa√m) is 1.0 × ^10⁻⁶ m· ^cycle⁻¹ or less.
[2] A casting step of casting a steel material having the component composition described in [1] at a casting speed of 1.8 m/min or less,
A heating process that involves heating to 1350°C or below,
A hot rolling process in which the steel material heated in the above heating process is rolled at a rolling completion temperature of 820°C or higher to form a steel pipe shape,
After the steel pipe obtained in the hot rolling process is held at a temperature of Ac ₃ or higher and 1000°C or lower, a cooling process is performed in which the cooling conditions are in group A or group B below.
A method for manufacturing a steel pipe, comprising a tempering step in which the steel pipe obtained in the cooling step is tempered at a temperature of 400°C or higher and an Ac value of ₁ or lower for less than 60 minutes.
Group A:
The average cooling rate from 800°C to 550°C is 15°C/s or more at the 1/4 wall thickness position from the inner surface of the steel pipe, and the average cooling rate from 550°C to 50°C is 15°C/s or less at the 1/4 wall thickness position from the inner surface of the steel pipe, cooling the steel pipe down to 50°C or below.
Group B:
The average cooling rate from 800°C to 300°C is 10°C/s or more at the 1/4 wall thickness position from the inner surface of the steel pipe, and the average cooling rate from 300°C to 50°C is 5°C/s or less at the 1/4 wall thickness position from the inner surface of the steel pipe, cooling the steel pipe down to 50°C or below.
[3] The method for manufacturing a steel pipe according to [2], further comprising a quenching step before the tempering step, in which the pipe is reheated to an Ac of ₃ or higher and 1000°C or lower, and the cooling conditions are those of group A or group B below.
Group A:
The average cooling rate from 800°C to 550°C is 15°C/s or more at the 1/4 wall thickness position from the inner surface of the steel pipe, and the average cooling rate from 550°C to 50°C is 15°C/s or less at the 1/4 wall thickness position from the inner surface of the steel pipe, cooling the steel pipe down to 50°C or below.
Group B:
The average cooling rate from 800°C to 300°C is 10°C/s or more at the 1/4 wall thickness position from the inner surface of the steel pipe, and the average cooling rate from 300°C to 50°C is 5°C/s or less at the 1/4 wall thickness position from the inner surface of the steel pipe, cooling the steel pipe down to 50°C or below.
[4] The method for manufacturing a steel pipe according to [2] or [3], wherein the casting speed is 1.0 m/min or less.
[5] In mass%,
C: 0.10-0.45%,
Si: 0.01-2.0%,
Mn: 0.3-2.0%,
Al: 0.01-0.15%,
N: 0.0005-0.008%,
P: 0.015% or less,
S: 0.0015% or less,
O: 0.01% or less,
H: 0.0010% or less,
Cu: 0 to 2.5%,
Ni: 0 to 2.5%,
Cr: 0-2.5%,
Mo: 0-2.0%,
Nb: 0 to 0.5%,
V: 0-0.5%,
Ti: 0 to 0.5%,
W: 0-2.5%,
B: 0 to 0.005%,
Sn: 0-0.3%,
Sb: 0 to 0.3%,
Ca: 0-0.01%,
Mg: 0 to 0.01%,
REM: 0~0.005%
It contains, with the remainder being Fe and unavoidable impurities, and has a component composition.
If the residual austenite content is 3% or less,
There are 20 or fewer carbides with a diameter of 200 nm or more per ¹⁰ μm².
A steel material exhibiting excellent fatigue properties in hydrogen, where the crack propagation rate da/dN in hydrogen at stress intensity of 1 MPa or higher (stress intensity factor range = 20 MPa√m) is 1.0 × ^10⁻⁶ m· ^cycle⁻¹ or less.
[6] A casting step of casting a steel material having the component composition described in [5] at a casting speed of 1.8 m/min or less,
A heating process that involves heating to 1350°C or below,
A hot rolling process is performed in which the steel material heated in the aforementioned heating process is rolled at a rolling completion temperature of 820°C or higher.
The steel material obtained in the hot rolling process is held at a temperature of Ac ₃ or higher and 1000°C or lower, followed by a cooling process where the cooling conditions are in group A or group B below.
A method for manufacturing steel, comprising: a tempering step in which the steel obtained in the cooling step is tempered at a temperature of 400°C or higher and an Ac of ₁ point or lower for less than 60 minutes.
Group A:
The average cooling rate from 800°C to 550°C is 15°C/s or more at a point 1/4 of the thickness from the surface of the steel material, and the average cooling rate from 550°C to 50°C is 15°C/s or less at a point 1/4 of the thickness from the surface of the steel material, cooling the steel material down to 50°C or below.
Group B:
The average cooling rate from 800°C to 300°C is 10°C/s or more at a point 1/4 of the thickness from the surface of the steel material, and the average cooling rate from 300°C to 50°C is 5°C/s or less at a point 1/4 of the thickness from the surface of the steel material, thereby cooling the steel material to 50°C or below.
[7] The method for manufacturing steel according to [6], further comprising a quenching step before the tempering step, in which the steel is reheated to an Ac of ₃ or higher and 1000°C or lower, and the cooling conditions are those of group A or group B below.
Group A:
The average cooling rate from 800°C to 550°C is 15°C/s or more at a point 1/4 of the thickness from the surface of the steel material, and the average cooling rate from 550°C to 50°C is 15°C/s or less at a point 1/4 of the thickness from the surface of the steel material, cooling the steel material down to 50°C or below.
Group B:
The average cooling rate from 800°C to 300°C is 10°C/s or more at a point 1/4 of the thickness from the surface of the steel material, and the average cooling rate from 300°C to 50°C is 5°C/s or less at a point 1/4 of the thickness from the surface of the steel material, thereby cooling the steel material to 50°C or below.
[8] The method for manufacturing steel according to [6] or [7], wherein the casting speed is 1.0 m/min or less.

According to the present invention, steel pipes and steel materials with extremely excellent fatigue characteristics under a high-pressure hydrogen gas environment can be obtained, which are extremely useful in industry.

Next, a specific method for carrying out the present invention will be described.

As the first embodiment, a method for implementing steel pipes will be specifically described, followed by a second embodiment, a method for implementing steel materials will be specifically described.

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

C: 0.10-0.45%
Carbon (C) is an element necessary for increasing strength. Below 0.10%, its effect is insufficient. Therefore, the C content should be 0.10% or more. Preferably, the C content should be 0.13% or more. More preferably, the C content should be 0.15% or more, and even more preferably, 0.18% or more. On the other hand, if it exceeds 0.45%, quench cracking may occur during hardening, and it can also cause the formation of coarse carbides, degrading the fatigue properties in hydrogen. Therefore, the C content should be 0.45% or less. Preferably, the C content should be 0.43% or less. More preferably, the C content should be 0.40% or less, and even more preferably, 0.38% or less.

Si: 0.01~2.0%
Si is included as a deoxidizing agent and an element that ensures hardenability during the steelmaking process, but its effect is insufficient if it is less than 0.01%, so the Si content should be 0.01% or more. Preferably, the Si content should be 0.1% or more. More preferably, the Si content should be 0.15% or more. On the other hand, if it exceeds 2.0%, the grain boundaries become brittle, degrading the low-temperature toughness and fatigue properties in hydrogen. Therefore, the Si content should be 2.0% or less. Preferably, the Si content should be 1.5% or less. Preferably, the Si content should be 1.0% or less, and more preferably 0.8% or less.

Mn: 0.3-2.0%
Mn is included as an element to ensure hardenability, but its effect is insufficient if it is less than 0.3%, so the Mn content should be 0.3% or more. A Mn content of 0.4% or more is preferable. A Mn content of 0.5% or more is more preferable. A Mn content of 0.6% or more is even preferable. On the other hand, if the content exceeds 2.0%, the grain boundary strength decreases and the low-temperature toughness deteriorates. Furthermore, a high Mn content increases austenite stability, which may exceed the specified amount of retained austenite and increase the amount of hydrogen in the steel. In addition, the hardness of the surface layer and the central segregation area increases during post-hot-rolling cooling (accelerated cooling or quenching), which deteriorates the fatigue properties in hydrogen. Therefore, the Mn content should be 2.0% or less. A Mn content of 1.5% or less is more preferable, and a Mn content of 1.3% or less is even preferable. A Mn content of 1.0% or less is most preferable.

Al: 0.01~0.15%
Al is included as a deoxidizing agent and, as fine precipitates of Al-based nitride, pinns austenite grains during heating, suppressing grain coarsening. However, this effect is insufficient if the Al content is less than 0.01%. Therefore, the Al content should be 0.01% or more. Preferably, the Al content should be 0.02% or more. More preferably, the Al content should be 0.03% or more. On the other hand, if the Al content exceeds 0.15%, the cleanliness of the steel decreases, and the toughness and fatigue properties in hydrogen deteriorate. Therefore, the Al content should be 0.15% or less. Preferably, the Al content should be 0.13% or less. More preferably, the Al content should be 0.10% or less, and even more preferably 0.08% or less.

N: 0.0005-0.008%
N is included because it forms fine precipitates by forming nitrides with Nb, Ti, Al, etc., and pinning the austenite grains during heating, thereby suppressing grain coarsening and improving low-temperature toughness. Since a content of less than 0.0005% does not sufficiently refine the structure, the N content should be 0.0005% or more. A N content of 0.001% or more is preferred. A N content of 0.0025% or more is more preferred. On the other hand, a content exceeding 0.008% increases the amount of solid-solution N, impairing the toughness of the base material and the heat-affected zone of the weld, and degrading the fatigue properties in hydrogen. Therefore, the N content should be 0.008% or less. A N content of 0.007% or less is preferred. A N content of 0.006% or less is more preferred, and 0.005% or less is even more preferred.

P: 0.015% or less. As an impurity element, P tends to segregate at grain boundaries, and if it exceeds 0.015%, it reduces the bonding strength between adjacent grains and degrades low-temperature toughness and fatigue properties in hydrogen. Therefore, the P content should be 0.015% or less. A P content of 0.013% or less is preferable, and a P content of 0.010% or less is more preferable. There is no particular lower limit, but it is preferable to set it at 0.001% or more to avoid increased costs.

S: 0.0015% or less. S, an impurity element, tends to segregate at grain boundaries and easily forms MnS, a nonmetallic inclusion. If it exceeds 0.0015%, the bonding strength between adjacent grains decreases, the amount of inclusions increases, and the low-temperature toughness and fatigue properties in hydrogen deteriorate. Therefore, the S content should be 0.0015% or less. A S content of 0.0013% or less is preferable. A S content of 0.0010% or less is more preferable, and 0.0008% or less is even more preferable. There is no particular lower limit, but it is preferable to set it at 0.0001% or more as this leads to increased costs.

O: 0.01% or less. Since oxygen forms oxides with Al and other elements, affecting the processability of the material, a lower amount is preferable. Content exceeding 0.01% increases inclusions and impairs processability. Furthermore, the fatigue properties in hydrogen deteriorate with increasing inclusions. Therefore, the oxygen content should be 0.01% or less. An oxygen content of 0.009% or less is preferable. An oxygen content of 0.008% or less is more preferable. There is no particular lower limit, but it is preferable to have a lower limit of 0.0001% or more as this leads to increased costs. An oxygen content of 0.002% or more is more preferable.

H: 0.0010% or less. Hydrogen 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 accelerates. Furthermore, since a high amount introduced increases the crack propagation rate, it is important to reduce the amount of hydrogen in the steel material. These effects are not a problem if the amount is 0.0010% or less, so the H content should be 0.0010% or less. Preferably, it is 0.0005% or less. More preferably, the H content is 0.0002% or less. On the other hand, if it is less than 0.00001%, it will be a factor in increasing costs, so it is preferable to have 0.00001% or more. It is preferable that the H content be 0.0001% or more. Note that the amount of hydrogen is the residual hydrogen after forming of the steel material, steel pipe, UOE, etc.

In the present invention, it is preferable that the steel composition consists of Fe and unavoidable impurities as the remainder of the above component composition. However, depending on the desired properties, it is preferable to further include one or more of the following, individually or simultaneously, as appropriate: Cu: 0-2.5%, Ni: 0-2.5%, Cr: 0-2.5%, Mo: 0-2.0%, Nb: 0-0.5%, V: 0-0.5%, Ti: 0-0.5%, W: 0-2.5%, B: 0-0.005%, Sn: 0-0.3%, Sb: 0-0.3%, Ca: 0-0.01%, Mg: 0-0.01%, and REM: 0-0.005%.

Cu: 0-2.5%
Cu has the effect of improving hardenability. For this reason, when Cu is included, the Cu content may be 0% or more, but since the above effect is difficult to obtain if it is less than 0.05%, it is preferable that the Cu content be 0.05% or more. On the other hand, if it exceeds 2.5%, it makes it easier for cracks to occur when the steel billet is heated while hot. For this reason, when Cu is included, it should be 2.5% or less. A Cu content of 2.3% or less is preferable. A Cu content of 2.0% or less is more preferable, and 1.8% or less is even more preferable.

Ni: 0-2.5%
Ni, like Cu, has the effect of improving hardenability and also improves toughness. For this reason, when Ni is included, the Ni content may be 0% or more, but since the above effects are difficult to obtain if it is less than 0.05%, it is preferable that the Ni content be 0.05% or more. On the other hand, if it exceeds 2.5%, it is not economically viable. Therefore, when Ni is included, it should be 2.5% or less. A Ni content of 2.3% or less is preferable. A content of 2.0% or less is more preferable, and 1.8% or less is preferable.

Cr: 0-2.5%
Cr is included as an element to ensure hardenability. When Cr is included, the Cr content may be 0% or more, but if it is less than 0.1%, the above effects are difficult to obtain, so it is preferable that the Cr content be 0.1% or more. On the other hand, if the content exceeds 2.5%, the toughness deteriorates and the economic efficiency is poor. Therefore, when Cr is included, it should be 2.5% or less. A Cr content of 2.3% or less is preferable. A Cr content of 2.0% or less is more preferable, 1.8% or less is even more preferable, and 1.5% or less is most preferable.

Mo: 0-2.0%
Mo has the effect of improving hardenability, so if Mo is included, the Mo content may be 0% or more, but if it is less than 0.05%, the above effect is difficult to obtain, so it is preferable that the Mo content be 0.05% or more. On the other hand, a content exceeding 2.0% is economically unfeasible. Therefore, if Mo is included, it should be 2.0% or less. A Mo content of 1.8% or less is preferable. A Mo content of 1.5% or less is more preferable, and 1.2% or less is even more preferable.

Nb: 0-0.5%
Nb has the effect of improving hardenability and, as fine precipitates of Nb-based carbonitrides, pinns austenite grains during heating, suppressing grain coarsening. For this reason, when Nb is included, the Nb content may be 0% or more, but since the above effect is difficult to obtain if it is less than 0.005%, it is preferable that the Nb content be 0.005% or more. It is even more preferable that the Nb content be 0.01% or more. On the other hand, if the content exceeds 0.5%, coarse Nb carbonitrides may precipitate, leading to a deterioration of toughness. Therefore, when Nb is included, the Nb content should be 0.5% or less. It is preferable that the Nb content be 0.4% or less. It is preferable that the Nb content be 0.3% or less, and it is preferable that the Nb content be 0.2% or less.

V: 0-0.5%
V has the effect of improving hardenability and, as fine precipitates of V-based carbides, pinns austenite grains during heating, suppressing grain coarsening. For this reason, when V is included, the V content may be 0% or more, but the above effect is difficult to obtain if it is less than 0.005%, so it is preferable that the V content be 0.005% or more. On the other hand, if the content exceeds 0.5%, coarse V carbonitrides may precipitate, leading to a deterioration of toughness. Therefore, when V is included, the V content should be 0.5% or less. A V content of 0.4% or less is preferable. A V content of 0.3% or less is more preferable, and 0.2% or less is even more preferable.

Ti: 0-0.5%
Ti has the effect of improving hardenability and, as fine precipitates of Ti-based carbonitrides, pinns austenite grains during heating, thereby suppressing grain growth. For this reason, when Ti is included, the Ti content may be 0% or more, but since the above effect is difficult to obtain if it is less than 0.005%, it is preferable that the Ti content be 0.005% or more. It is preferable that the Ti content be 0.01% or more. On the other hand, if the content exceeds 0.5%, coarse angular nitrides tend to form, and toughness deteriorates. Therefore, when Ti is included, the Ti content should be 0.5% or less. It is preferable that the Ti content be 0.4% or less. It is more preferable that the Ti content be 0.3% or less, and even more preferable that it be 0.2% or less.

W: 0-2.5%
Since W has the effect of improving hardenability, if W is included, the W content may be 0% or more, but if it is less than 0.05%, the above effect is difficult to obtain, so it is preferable that the W content be 0.05% or more. On the other hand, if it exceeds 2.5%, it is not economically viable. Therefore, if W is included, the W content should be 2.5% or less. A W content of 2.3% or less is preferable. A W content of 2.0% or less is more preferable, and 1.8% or less is even more preferable.

B: 0-0.005%
Since B is an element that ensures hardenability, if B is included, the B content may be 0% or more, but if it is less than 0.0005%, the above effect is difficult to obtain, so it is preferable that the B content be 0.0005% or more. On the other hand, if it exceeds 0.005%, the toughness deteriorates. Therefore, if B is included, the B content should be 0.005% or less. It is preferable that the B content be 0.004% or less. It is more preferable that the B content be 0.003% or less, and even more preferable that it be 0.002% or less.

Sn: 0-0.3%
Sn has the effect of improving the corrosion resistance of steel pipes. For this reason, when Sn is included, the Sn content may be 0% or more, but since the above effect is difficult to obtain if it is less than 0.005%, it is preferable that the Sn content be 0.005% or more. It is more preferable that the Sn content be 0.01% or more. On the other hand, if the content exceeds 0.3%, the high-temperature ductility decreases and the possibility of cracking during casting increases. For this reason, when Sn is included, the Sn content should be 0.3% or less. It is preferable that the Sn content be 0.25% or less. It is more preferable that the Sn content be 0.2% or less, and even more preferable that be 0.15% or less.

Sb: 0-0.3%
Sb has the effect of improving the corrosion resistance of steel pipes. For this reason, when Sb is included, the Sb content may be 0% or more, but since the above effect is difficult to obtain if it is less than 0.005%, it is preferable that the Sb content be 0.005% or more. It is more preferable that the Sb content be 0.01% or more. On the other hand, if the content exceeds 0.3%, the high-temperature ductility decreases and the hot-rolling ability decreases. For this reason, when Sb is included, the Sb content should be 0.3% or less. It is preferable that the Sb content be 0.25% or less. It is more preferable that the Sb content be 0.2% or less, and even more preferable that be 0.15% or less.

Ca: 0-0.01%
Ca has the effect of controlling the morphology of sulfide-based inclusions, transforming them into CaS, a spherical inclusion that is difficult to stretch by rolling, instead of MnS, an inclusion that is easily stretched by rolling. For this reason, when Ca is included, the Ca content may be 0% or more, but the above effect is difficult to obtain if it is less than 0.0005%, so it is preferable that the Ca content be 0.0005% or more. It is more preferable that the Ca content be 0.001% or more. On the other hand, if the content exceeds 0.01%, the cleanliness decreases, and the material properties such as toughness deteriorate. Therefore, when Ca is included, the Ca content should be 0.01% or less. It is preferable that the Ca content be 0.005% or less. It is more preferable that the Ca content be 0.003% or less, and even more preferable that it be 0.002% or less.

Mg: 0-0.01%
Mg may be used as a desulfurizing agent for molten iron. Therefore, if Mg is included, the Mg content may be 0% or more, but since the above effects are difficult to obtain if it is less than 0.0005%, it is preferable that the Mg content be 0.0005% or more. It is more preferable that the Mg content be 0.001% or more. On the other hand, a content exceeding 0.01% leads to a decrease in cleanliness. Therefore, if Mg is included, the Mg content should be 0.01% or less. It is preferable that the Mg content be 0.005% or less. It is more preferable that the Mg content be 0.004% or less, and even more preferable that it be 0.003% or less.

REM: 0~0.005%
REM improves resistance to SR cracking by reducing the amount of dissolved sulfur (S) at grain boundaries in steel by forming sulfides as REM(O,S). Therefore, when REM is included, the REM content may be 0% or more, but since the above effect is difficult to obtain if it is less than 0.0005%, it is preferable that the REM content be 0.0005% or more. On the other hand, if the content exceeds 0.005%, REM sulfides accumulate significantly in the precipitated crystal zone, leading to deterioration of the material. Therefore, when REM is included, the REM content should be 0.005% or less. A REM content of 0.003% or less is preferable. A REM content of 0.001% or less is even more preferable. REM stands for Rare Earth Metal, and it is a rare earth metal.

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 preferred metallic structure of the steel pipe of the present invention will be described in detail.

Retained austenite is 3% or less. The presence of retained austenite in steel pipes can increase the amount of hydrogen in the steel, potentially increasing its susceptibility to hydrogen embrittlement. Furthermore, if austenite transforms into martensite due to stress loading during use, martensite is very hard and therefore prone to hydrogen cracking, which can cause cracks to initiate from the martensite portion. In this invention, the fatigue crack propagation rate is reduced by limiting the retained austenite to 3% or less. Preferably, it is 2% or less, and more preferably 1% or less. The retained austenite may be 0%.

In the steel pipe of the present invention, if the amount of coarse carbides exceeds a predetermined value, it adversely affects the fatigue characteristics in hydrogen. For this reason, the amount of carbides with a diameter of 200 nm or more is set to 20 ^or less per ¹⁰ μm². Preferably, it is 15 or less per ¹⁰ μm². More preferably, it is 10 or less per ¹⁰ μm², and even more preferably, 5 or less per ¹⁰ μm². For the lower limit, a smaller number is better, and it may be 0 or less per ¹⁰ μm². Here, the diameter refers to the value calculated from 2√(A/2 × B/ ₂ ) using the long side A and short side B passing through the center. Carbides refer to intermetallic compounds including, for example, cementite, ε-carbide, χ-carbide, and _Fe7C3 .

Furthermore, when carbides with a diameter of less than 200 nm are uniformly dispersed within the grain at intervals of 100 nm or more, they do not adversely affect the fatigue properties in hydrogen and contribute to improved strength. Therefore, it is preferable that the amount of carbides with a diameter of less than 200 nm dispersed within the grain be 10 or more per ¹⁰ μm². As for the upper limit, it is preferable that the amount of carbides with a diameter of less than 200 nm dispersed within the grain be 100 or less per ¹⁰ μm² because too many precipitates lead to coarsening.
Furthermore, since carbide precipitation is influenced by the conditions of the cooling and tempering processes after hot rolling, which will be described later, it is important to control these conditions.

The stress intensity factor in hydrogen at 1 MPa or higher is such that the crack propagation rate da/dN at 20 MPa√m is 1.0 × ^10⁻⁶ m· ^cycle⁻¹ or less. Fatigue crack propagation rate is an important parameter in the design of steel pipes used in line pipes and gas containers, and is necessary to ensure the safety of structural members and obtain a sufficient service life. In structural members, it is difficult to eliminate cracks or crack initiation sites entirely, and cracks inevitably occur and propagate when subjected to repeated stress. The crack propagation rate is small when the stress state at the crack tip is small, and increases as the stress state at the crack tip increases. In a hydrogen environment, hydrogen penetrates into the steel pipe, making it easier for cracks to propagate. The degree to which hydrogen accelerates crack propagation rate is greatly influenced by the material's microstructure and precipitates. In crack propagation tests in hydrogen at pressures of 1 MPa or higher, if the crack propagation rate da/dN in the stress intensity factor range = 20 MPa√m, obtained by conducting fatigue tests in accordance with ASTM E647 with a frequency of 1 Hz, repetition waveform: sine wave, control method: load control, and stress ratio: R = 0.1, is 1.0 × ^10⁻⁶ m· ^cycle⁻¹ or less, then the service life of steel structures in a high-pressure hydrogen environment can be sufficiently ensured. Therefore, in crack propagation tests in hydrogen at pressures of 1 MPa or higher, the crack propagation rate da/dN in the stress intensity factor range = 20 MPa√m is set to 1.0 × ^10⁻⁶ m· ^cycle⁻¹ or less. The crack propagation rate da/dN at a stress intensity factor of 20 MPa√m is preferably 0.9 × ^10⁻⁶ m· ^cycle⁻¹ or less, more preferably 0.8 × ^10⁻⁶ m· ^cycle⁻¹ or less, and even more preferably 0.7 × ^10⁻⁶ m· ^cycle⁻¹ or less. For the lower limit, it is considered better if it is close to the result in air, and it is more preferable that the crack propagation rate da/dN at a stress intensity factor of 20 MPa√m is preferably 0.05 × ^10⁻⁶ m· ^cycle⁻¹ or more.

Furthermore, while the thickness of the steel pipe is not particularly limited, a thickness of 5 mm or more is preferred. A thickness of 30 mm or less is preferred.

The steel pipes of this invention include seamless steel pipes, electric resistance welded pipes, UOE steel pipes, and the like. Below, a method for manufacturing a seamless steel pipe will be specifically described as an example.

The steel pipe exhibiting excellent fatigue characteristics in hydrogen gas according to the present invention has the above-mentioned component composition and only requires that the crack propagation rate in hydrogen is met. The manufacturing method thereof will be specifically described below.

It goes without saying that by performing processing to achieve a similar thermal history, it is possible to manufacture electric resistance welded pipes and UOE steel pipes.

The steel pipe of the present invention can be manufactured by sequentially carrying out the following steps (1) to (3).
(1) A process of casting steel material after adjusting its composition. (2) A process of hot rolling and cooling (accelerated cooling) of the cast material, in which the material is heated and rolled to obtain the shape of a steel pipe (including cases where the material is reheated and hardened before the tempering process).
(3) A step of tempering the steel pipe obtained in the above step.

The following describes each process. Unless otherwise specified, the temperature in the following description refers to the temperature at the center of the steel material or steel pipe's thickness. 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 consider the heat transfer coefficient of the steel material, based on the steel pipe surface temperature measured with a radiation thermometer.

[Casting Process]
Casting speed: 1.8 m/min or less. The slower the casting speed, the more the hydrogen concentration and inclusions in the steel can be reduced, and this effect becomes significant at 1.8 m/min or less. Therefore, the casting speed should be 1.8 m/min or less. Preferably, it is 1.5 m/min or less. More preferably, it is 1.0 m/min or less. Even more preferably, it is 0.5 m/min or less. Most preferably, it is 0.1 m/min or less. There is no particular lower limit, but the casting speed should be greater than 0 m/min.

Heating Process: In order to perform hot rolling, the steel material having the above-described component composition is heated. The steel material is not particularly limited, but for example, slabs or billets obtained by a conventional continuous casting method can be used.

Heating temperature: 1350°C or lower. If the heating temperature in the heating process exceeds 1350°C, the average particle size of the prior austenite grains becomes excessively large, and various properties deteriorate. Therefore, the heating temperature should be 1350°C or lower. It is more preferable that the heating temperature be 1300°C or lower, even more preferable that it be 1250°C or lower, and most preferable that it be 1200°C or lower. On the other hand, a lower heating temperature is preferable because it reduces the amount of hydrogen in the steel, but if it is too low, the finish rolling temperature will decrease, making rolling difficult. Therefore, it is preferable that the heating temperature be 950°C or higher. It is more preferable that the heating temperature be 1000°C or higher. There is no particular limit to the heating time, but if it is too long, the risk of an increase in hydrogen introduced into the steel pipe increases, so it is preferable to keep it at 180 minutes or less. It is more preferable that the heating time be 150 minutes or less, and even more preferable that it be 120 minutes or less. There is no particular lower limit, but it is preferable that the heating time be 30 minutes or more, and more preferable that it be 60 minutes or more.

Rolling Process The steel material heated in the above heating process is rolled under the following conditions to form a steel pipe shape. For the rolling process, a standard Mannesmann-plug mill method or a Mannesmann-mandrel mill method, including perforated rolling, can be used for hot rolling.

Rolling end temperature: 820°C or higher. If the rolling end temperature is below 820°C, the rolling load will be excessive, increasing the risk of rolling problems. Therefore, the rolling end temperature should be 820°C or higher. Preferably, the rolling end temperature should be 850°C or higher, and more preferably 900°C or higher. On the other hand, there is no particular upper limit to the rolling end temperature, but if the temperature is too high, the metal structure tends to become non-uniform, so it is preferable that the rolling end temperature be 1200°C or lower. More preferably, the rolling end temperature should be 1150°C or lower, and even more preferably 1100°C or lower.

[Cooling process (accelerated cooling process)]
In the cooling process, the steel material having the above-described component composition is heated and held to a temperature of Ac ₃ or higher and 1000°C or lower, either as is or after being processed into a steel pipe, and then cooled under the cooling conditions of Group A or Group B below. It is preferable to hold the material at the above temperature for 10 minutes or more. It is more preferable to hold it for 15 minutes or more, and even more preferable to hold it for 20 minutes or more. There is no particular upper limit, but it is preferable to hold the material at the above temperature for 60 minutes or less, and more preferable to hold it for 45 minutes or less.

Heating temperature after steel pipe processing: Ac ₃ point or higher and 1000°C or lower. If the heating temperature in the cooling process is less than Ac ₃ point, ferrite will remain in the steel after cooling, reducing the strength of the steel pipe and its fatigue properties in hydrogen. Therefore, the heating temperature should be Ac ₃ point or higher. Preferably, the heating temperature should be Ac ₃ point + 30°C or higher, and more preferably, Ac ₃ point + 50°C or higher. However, for composition systems where Ac ₃ point + 30°C and Ac ₃ point + 50°C exceed 1000°C, the above Ac ₃ point + 30°C and Ac ₃ point + 50°C do not apply. On the other hand, if the heating temperature is higher than 1000°C, the austenite grains may coarseen, which may cause a decrease in the impact absorption energy value and toughness of the material after heat treatment. Therefore, the heating temperature should be 1000°C or lower. Preferably, the heating temperature should be 950°C or lower, and more preferably, 900°C or lower. However, for component systems where the Ac score is less than ₃ at 950°C and 900°C, the above-mentioned limits of 950°C and 900°C do not apply.
In this cooling process, if the temperature after rolling satisfies the heating conditions, the material may be cooled directly, or it may be reheated and cooled after rolling. Furthermore, if the steel sheet has been cooled by air cooling, it may be reheated to a temperature of Ac ₃ or higher but 1000°C or lower, and then cooled under the cooling conditions of Group A or Group B below.
In this invention, the Ac ₃ point (°C) is calculated using the following formula.
Ac ₃ (℃)=910-203[C]1/2-30[Mn]+44.7[Si]+700[P]+100[Al]+31.5[Mo]-11[Cr]-15.2[Ni]-20[Cu]+104[V]
However, [M] in the formula represents the content (mass%) of element M.

Average cooling rate Group A: Cooling to 50°C or below under the conditions that the average cooling rate in the range from 800°C to 550°C at the 1/4 wall thickness position from the inner surface of the steel pipe is 15°C/s or more, and the average cooling rate from 550°C to 50°C is 15°C/s or less. If the average cooling rate from 800°C to 550°C at the 1/4 wall thickness position from the inner surface of the steel pipe is less than 15°C/s, the predetermined carbide density cannot be obtained. Furthermore, although the microstructure is not particularly limited, it is preferable to obtain 90% or more of either bainite or martensite by area ratio in order to obtain the predetermined fatigue characteristics in hydrogen. If the above average cooling rate is less than 15°C/s, it becomes difficult to obtain 90% or more of bainite by area ratio, but depending on the component composition, it may also affect the formation of martensite. For this reason, the average cooling rate at the 1/4 wall thickness position from the inner surface of the steel pipe should be 15°C/s or more. From the viewpoint of suppressing variations in microstructure, it is preferable that the average cooling rate be 17°C/s or more. The average cooling rate from 800°C to 550°C is more preferably 20°C/s or higher, and most preferably 22°C/s or higher. On the other hand, in order to suppress variations in particle size, the average cooling rate is preferably 50°C/s or lower, more preferably 45°C/s or lower, and even more preferably 40°C/s or lower.
Furthermore, by cooling to 50°C or below under conditions where the average cooling rate from 550°C to 50°C is 15°C/s or less, retained austenite can be reduced, thereby reducing the amount of hydrogen in the steel. For this reason, the average cooling rate from 550°C to 50°C should be 15°C/s or less. It is more preferable that the average cooling rate from 550°C to 50°C be 12°C/s or less, and even more preferable that it be 10°C/s or less. There is no particular lower limit, but it is preferable that the average cooling rate from 550°C to 50°C be 1°C/s or more.
The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling can be used alone or in combination, but water cooling or oil cooling is preferred from 800°C to 550°C, and air cooling is preferred from 550°C to 50°C.

Group B: Cooling to 50°C or below under the conditions that the average cooling rate from 800°C to 300°C at the 1/4 wall thickness position from the inner surface of the steel pipe is 10°C/s or more, and the average cooling rate from 300°C to 50°C is 5°C/s or less. If the average cooling rate from 800°C to 300°C at the 1/4 wall thickness position from the inner surface of the steel pipe is less than 10°C/s, the predetermined carbide density cannot be obtained, and deterioration of fatigue properties occurs. Also, if the above average cooling rate is less than 10°C/s, it becomes difficult to obtain a martensite area ratio of 90% or more, but depending on the component composition, it may affect the formation of bainite. For this reason, the average cooling rate at the 1/4 wall thickness position from the inner surface of the steel pipe should be 10°C/s or more. From the viewpoint of suppressing variations in microstructure, it is preferable that the average cooling rate from 800°C to 300°C be 12°C/s or more, more preferably 15°C/s or more, and even more preferably 17°C/s or more. While there is no particular upper limit, it is preferable that the average cooling rate be 60°C/s or less.
Furthermore, the amount of hydrogen in the steel can be reduced by cooling it to 50°C or below under conditions where the average cooling rate from 300°C to 50°C is 5°C/s or less. For this reason, the average cooling rate from 300°C to 50°C should be 5°C/s or less. Preferably, the average cooling rate from 300°C to 50°C should be 1°C/s or less. There is no particular limit to the lower limit, but it is preferable to have a lower limit of 0.1°C/s or more.
The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling can be used alone or in combination, but water cooling or oil cooling is preferred from 800°C to 300°C, and air cooling is preferred from 300°C to 50°C.

[Reheating and quenching process (preferred conditions)]
Reheating temperature before quenching: Ac ₃ point or higher and 1000°C or lower. If the temperature at the center of the plate thickness is below Ac ₃ point, some untransformed austenite will remain, making it impossible to obtain the desired steel structure after hot rolling, quenching, and tempering (described later). For this reason, the pre-quenching heating temperature during reheating should be Ac ₃ point or higher. Preferably, it should exceed Ac ₃ point. Furthermore, in order to suppress excessive coarsening of the initial austenite grain size and improve production efficiency, it is preferable that the pre-quenching heating temperature be 1000°C or lower. More preferably, it should be 980°C or lower, and even more preferably 960°C or lower. Most preferably, it should be 950°C or lower. By setting the pre-quenching reheating temperature to a lower temperature within the range of Ac ₃ point or higher, the initial austenite grain size can be refined, and the hydrogen crack propagation rate can be reduced.

Average cooling rate during quenching: Group A or Group B below Group A: Cooling to 50°C or below under the conditions that the average cooling rate in the range from 800°C to 550°C at the 1/4 wall thickness position from the inner surface of the steel pipe is 15°C/s or more, and the average cooling rate from 550°C to 50°C is 15°C/s or less. If the average cooling rate from 800°C to 550°C at the 1/4 wall thickness position from the inner surface of the steel pipe is less than 15°C/s, the predetermined carbide density cannot be obtained. Also, if the above average cooling rate is less than 15°C/s, it becomes difficult to obtain bainite in area ratio of 90% or more, but depending on the component composition, it may affect the formation of martensite. For this reason, the average cooling rate at the 1/4 wall thickness position from the inner surface of the steel pipe should be 15°C/s or more. From the viewpoint of suppressing variations in microstructure, it is preferable that the average cooling rate be 17°C/s or more, more preferably 20°C/s or more, and even more preferably 22°C/s or more. On the other hand, in order to suppress variations in particle size, the average cooling rate is preferably 50°C/s or less, more preferably 47°C/s or less, and even more preferably 45°C/s or less.
Furthermore, by cooling to 50°C or below under conditions where the average cooling rate from 550°C to 50°C is 15°C/s or less, retained austenite can be reduced, thereby reducing the amount of hydrogen in the steel. For this reason, the average cooling rate from 550°C to 50°C should be 15°C/s or less. Preferably, the average cooling rate from 550°C to 50°C should be 12°C/s or less, and more preferably 10°C/s or less. There is no particular lower limit, but preferably, the average cooling rate from 550°C to 50°C should be 1°C/s or more.
The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling can be used alone or in combination, but water cooling or oil cooling is preferred from 800°C to 550°C, and air cooling is preferred from 550°C to 50°C.

Group B: Cooling to 50°C or below under the conditions that the average cooling rate from 800°C to 300°C at the 1/4 wall thickness position from the inner surface of the steel pipe is 10°C/s or more, and the average cooling rate from 300°C to 50°C is 5°C/s or less. If the average cooling rate from 800°C to 300°C at the 1/4 wall thickness position from the inner surface of the steel pipe is less than 10°C/s, the predetermined carbide density cannot be obtained, and the fatigue properties deteriorate. Also, if the above average cooling rate is less than 10°C/s, it becomes difficult to obtain a martensite area ratio of 90% or more, but depending on the component composition, it may affect the formation of bainite. For this reason, the average cooling rate at the 1/4 wall thickness position from the inner surface of the steel pipe should be 10°C/s or more. From the viewpoint of suppressing variations in the structure, it is preferable that the average cooling rate be 17°C/s or more, more preferably 20°C/s or more, and even more preferably 25°C/s or more. On the other hand, although there is no particular upper limit specified for the average cooling rate, if it exceeds 60°C/s, a large amount of hard structure will be generated on the surface of the steel plate, making it impossible to obtain a steel structure having the structure intended in the present invention, and the fatigue properties in hydrogen will decrease. Therefore, it is preferable that the average cooling rate be 60°C/s or less.
Furthermore, the amount of hydrogen in the steel can be reduced by cooling it to 50°C or below under conditions where the average cooling rate from 300°C to 50°C is 5°C/s or less. For this reason, the average cooling rate from 300°C to 50°C is set to 5°C/s or less. Preferably, the average cooling rate is 3°C/s or less, and more preferably 1°C/s or less. There is no particular limit to the lower limit, but it is preferable to set it to 0.1°C/s or more.
The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling can be used alone or in combination, but water cooling or oil cooling is preferred from 800°C to 300°C, and air cooling is preferred from 300°C to 50°C.

Cooling stop temperature during quenching: 50°C or lower. If the cooling stop temperature exceeds 50°C, the desired carbide density cannot be obtained, and furthermore, the above transformation will not be completed, making it impossible to obtain the desired steel structure after tempering. For this reason, quenching should be performed to a temperature of 50°C or lower. The cooling stop temperature is preferably 45°C or lower, and more preferably 40°C or lower. There is no particular lower limit, but it is preferable that the cooling stop temperature be 25°C or higher.

Tempering Process Tempering temperature: 400°C or higher, below Ac ₁ point By raising the temperature under conditions where the average heating rate is 0.01°C/s or higher, and setting the tempering temperature to 400°C or higher, it is possible to reduce austenite and hydrogen in the steel, and obtain a predetermined carbide density. The tempering temperature is preferably 450°C or higher, and more preferably 500°C or higher. On the other hand, if the temperature is raised above Ac ₁ point, the amount of austenite and hydrogen in the steel may increase. For this reason, the tempering temperature should be below Ac ₁ point. Preferably, it is in the range of (Ac ₁ point - 30)°C or lower. The upper limit of the average heating rate during tempering is not particularly limited, but it is preferably 1°C/s or lower. If the tempering time is too long, the carbides will coarseen and adversely affect hydrogen embrittlement, so it should be less than 60 minutes. The tempering time is preferably 50 minutes or less. If the tempering time is too short, the austenite in the steel will not be reduced and the amount of hydrogen will not be reduced. Therefore, it is preferable to temper for 10 minutes or more, and more preferable to temper for 20 minutes or more.

In this invention, the method for determining Ac ₁ point (°C) is not specifically defined, but for example,
Ac ₁ can be calculated as Ac 1 = 723 - 14Mn + 22Si - 14.4Ni + 23.3Cr. In the above formula, each element symbol represents the mass %) of each element in the steel, and elements that are not present are represented as 0.

Dehydrogenation Treatment Process: When hydrogen is present in steel materials, 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 sheets with excellent fatigue resistance 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)
The hydrogen diffusion coefficient varies depending on the components and metal structure, but for example, a hydrogen diffusion coefficient of 1 × ^10⁻⁵ to 5 × ^{10⁻³ mm²} /s may be used. More preferably, it is 5 × ^{10⁻⁴ mm²} /s or less.
The dehydrogenation treatment process is carried out before pipe fabrication or 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. On the other hand, if the temperature 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. Furthermore, 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 costs. 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 mentioned here refers to the temperature of the atmosphere during the dehydrogenation treatment process. Room temperature is defined as 20 ± 10°C.

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 ambient temperature satisfies the holding time R (sec), 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 dehydrogenation treatment 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 T at the dehydrogenation treatment temperature 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 dehydrogenation treatment temperature T. In other words, at least the former method allows for appropriate control of the amount of hydrogen in the surface layer of steel materials and steel pipes, and if the latter method is also implemented, the amount of hydrogen in the steel materials and steel pipes from the surface layer to the center of the plate thickness can be appropriately controlled. The plate thickness temperature and the plate thickness center temperature Tc can be measured using thermocouples or the like, or predicted using the finite element method or the like.

Furthermore, since scale on the steel surface inhibits dehydrogenation, it is preferable to remove the scale before performing dehydrogenation treatment. The method of scale removal is not limited; for example, physical cleaning by high-pressure washing or chemical methods using scale removers may be used. The thickness of the scale removal is not limited, but removing approximately 100 μm is sufficient to achieve the desired scale removal effect.

Second Embodiment The steel material of the present invention will now be described in detail. The composition, microstructure, and crack propagation rate of the steel material are the same as those described for steel pipes, and the manufacturing method is also carried out in the same manner as described for steel pipes, except for the rolling and cooling processes (casting, heating, reheating/quenching, tempering, and dehydrogenation treatment). The rolling and cooling processes are carried out as follows.

Rolling Process The steel material heated in the heating process described above in the steel pipe manufacturing method is hot-rolled in a hot rolling mill under the following conditions.

Rolling end temperature: 820°C or higher. If the rolling end temperature is below 820°C, the rolling load will be excessive, increasing the risk of rolling problems. Therefore, the rolling end temperature should be 820°C or higher. Preferably, the rolling end temperature should be 850°C or higher, and more preferably 900°C or higher. On the other hand, there is no particular upper limit to the rolling end temperature, but if the temperature is too high, the metal structure tends to become non-uniform, so it is preferable that the rolling end temperature be 1200°C or lower. More preferably, the rolling end temperature should be 1150°C or lower, and even more preferably 1100°C or lower.

Cooling process (accelerated cooling process)
In the cooling process, after hot rolling the steel material having the above-described component composition, it is heated and held at a temperature of Ac ₃ or higher and 1000°C or lower, and then cooled under the cooling conditions of Group A or Group B below. It is preferable to hold it at the above temperature for 10 minutes or more. It is more preferable to hold it for 15 minutes or more, and even more preferable to hold it for 20 minutes or more. There is no particular upper limit, but it is preferable to hold it at the above temperature for 60 minutes or less, and more preferable to hold it for 45 minutes or less.

Heating temperature after hot rolling: Ac ₃ point or higher and 1000°C or lower. If the heating temperature in the cooling process is less than Ac ₃ point, ferrite will remain in the steel after cooling, reducing the strength and fatigue properties of the steel. Therefore, the heating temperature should be Ac ₃ point or higher. Preferably, the heating temperature should be Ac ₃ point + 30°C or higher, and more preferably Ac ₃ point + 50°C or higher. However, for composition systems where Ac ₃ point + 30°C and Ac ₃ point + 50°C exceed 1000°C, the above Ac ₃ point + 30°C and Ac ₃ point + 50°C do not apply. On the other hand, if the heating temperature is higher than 1000°C, the austenite grains may coarseen, which may cause a decrease in the impact absorption energy value and toughness of the material after heat treatment. Therefore, the heating temperature should be 1000°C or lower. More preferably, it should be 950°C or lower, and even more preferably 900°C or lower. However, for component systems where the Ac score is less than ₃ at 950°C and 900°C, the above-mentioned limits of 950°C and 900°C do not apply.
In this cooling process, if the temperature after rolling satisfies the heating conditions, the material may be cooled directly, or it may be reheated and cooled after rolling. Furthermore, if the steel material has been cooled by air cooling, it may be reheated to a temperature of Ac ₃ or higher but 1000°C or lower, and then cooled under the cooling conditions of Group A or Group B below (this is called quenching).
In this invention, the Ac ₃ point (°C) is calculated using the following formula.
Ac ₃ (℃)=910-203[C]1/2-30[Mn]+44.7[Si]+700[P]+100[Al]+31.5[Mo]-11[Cr]-15.2[Ni]-20[Cu]+104[V]
However, [M] in the formula represents the content (mass%) of element M.

Average cooling rate Group A: Cooling to 50°C or below under the conditions that the average cooling rate in the range from 800°C to 550°C at a position 1/4 of the plate thickness from the steel surface is 15°C/s or more, and the average cooling rate from 550°C to 50°C is 15°C/s or less. If the average cooling rate from 800°C to 550°C at a position 1/4 of the plate thickness from the steel surface is less than 15°C/s, the predetermined carbide density cannot be obtained. Furthermore, although the microstructure is not particularly limited, it is preferable to obtain 90% or more of either bainite or martensite by area ratio in order to obtain predetermined fatigue characteristics in hydrogen. If the above average cooling rate is less than 15°C/s, it becomes difficult to obtain 90% or more of bainite by area ratio, but depending on the component composition, it may also affect the formation of martensite. For this reason, the average cooling rate at a position 1/4 of the plate thickness from the steel surface should be 15°C/s or more. From the viewpoint of suppressing variations in microstructure, it is preferable that the average cooling rate be 17°C/s or more. It is more preferable to have a cooling rate of 20°C/s or higher, and even more preferable to have a cooling rate of 22°C/s or higher. On the other hand, in order to suppress variations in particle size, the average cooling rate should be 50°C/s or lower. It is more preferable to have a cooling rate of 47°C/s or lower, and even more preferable to have a cooling rate of 45°C/s or lower.
Furthermore, by cooling the steel to 50°C or below under conditions where the average cooling rate from 550°C to 50°C is 15°C/s or less, retained austenite can be reduced, thereby reducing the amount of hydrogen in the steel. For this reason, the average cooling rate from 550°C to 50°C should be 15°C/s or less. The lower limit is not particularly limited, but it is preferable that the average cooling rate from 550°C to 50°C be 1°C/s or more.
The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling can be used alone or in combination, but water cooling or oil cooling is preferred from 800°C to 550°C, and air cooling is preferred from 550°C to 50°C.

Group B: Cooling to 50°C or below under the conditions that the average cooling rate from 800°C to 300°C at a position 1/4 of the plate thickness from the steel surface is 10°C/s or more, and the average cooling rate from 300°C to 50°C is 5°C/s or less. If the average cooling rate from 800°C to 300°C at a position 1/4 of the plate thickness from the steel surface is less than 10°C/s, the predetermined carbide density cannot be obtained, and deterioration of fatigue properties occurs. Also, if the above average cooling rate is less than 10°C/s, it becomes difficult to obtain a martensite area ratio of 90% or more, but depending on the component composition, it may affect the formation of bainite. For this reason, the average cooling rate at a position 1/4 of the plate thickness from the steel surface should be 10°C/s or more. From the viewpoint of suppressing variations in the structure, it is more preferable to set it to 12°C/s or more. It is more preferable for the average cooling rate to be 15°C/s or more, and even more preferable to be 17°C/s or more. On the other hand, although there is no particular upper limit specified for the average cooling rate, if it exceeds 60°C/s, a large amount of hard structure will be generated on the surface of the steel plate, making it impossible to obtain a steel structure having the structure intended for this invention, and the fatigue properties in hydrogen will decrease. Therefore, it is preferable that the average cooling rate be 60°C/s or less.
Furthermore, the amount of hydrogen in the steel can be reduced by cooling it to 50°C or below under the condition that the average cooling rate from 300°C to 50°C is 5°C/s or less. For this reason, the average cooling rate from 300°C to 50°C is set to 5°C/s or less. Preferably, the average cooling rate is 1°C/s or less, and more preferably 0.8°C/s or less. There is no particular limit to the lower limit, but it is preferable to set it to 0.1°C/s or more.
The cooling method is not particularly limited, and any method such as water cooling, oil cooling, or air cooling can be used alone or in combination, but water cooling or oil cooling is preferred from 800°C to 300°C, and air cooling is preferred from 300°C to 50°C.

While this is not necessary for thick plates, it is preferable to wind thin steel plates into a coil.

Furthermore, the steel material exhibiting excellent fatigue properties in hydrogen gas according to the present invention includes various classifications such as thin plates, thick plates, and steel pipes having the above-mentioned component composition and exhibiting excellent fatigue crack propagation characteristics in hydrogen gas, or it may be a steel material for hydrogen pipelines formed into a predetermined shape.

Under the above conditions, steel pipes and steel materials with excellent fatigue properties in hydrogen that satisfy a predetermined crack propagation rate in hydrogen can be obtained.

The following describes examples that verify the effects of the present invention. The following are preferred examples of the present invention, and the present invention is not limited in any way by these examples. In the following examples, the manufacturing method and characteristic evaluation of seamless steel pipes for actual steel structures were investigated.

Billets with the component compositions shown for steel pipes No. 1 to 29 (billet No. A to AC) and 40 to 87 (billet No. AN to CI) in Tables 1-1 and 1-2 were manufactured at a casting speed of 0.6 m/min. The billets were heated to 1250°C and expanded to obtain seamless steel pipes. The steel pipes were manufactured under conditions that the expansion was completed at 820°C or higher. For steel pipes with an _Ac3 point of 950°C or lower, the obtained steel pipes were heated and held at 950°C, and for steel pipes with an _Ac3 point exceeding 950°C, they were heated and held at 1000°C. After that, they were water-cooled under the conditions described in Tables 2-1 and 2-2, and then tempered. The metallographic structure and mechanical properties were evaluated. Also, for steel materials No. 1 to 29 in Tables 1-1 and 1-2... Slabs with the component compositions shown in slabs 30-39 (slab No. AD-AM) were produced at a casting speed of 0.6 m/min, heated to 1250°C, and then rolled in a hot rolling mill at 820°C or higher. The resulting steel materials were heated and held at 950°C for those with an Ac3 _point of 950°C or lower, and heated and held at 1000°C for those with an Ac3 _point exceeding 950°C. After that, the resulting steel materials were water-cooled under the conditions described in Tables 2-1 and 2-2, then tempered, and their microstructure and mechanical properties were evaluated in the same way as for steel pipes. After tempering, dehydrogenation treatment was also performed on Nos. 2, 5, 14, 15, 43, and 63-69. For the dehydrogenation treatment, the dehydrogenation treatment temperature T, which is the ambient temperature, was held at 50°C for 3 hours, and then allowed to cool naturally. The evaluation method is as follows. The tempering temperature was arbitrarily adjusted so that the tensile strength of the material was in the range of 520 to 700 MPa.

Furthermore, billets with the component compositions shown for steel pipes No. 88 to 101 (billet No. AO1 to BB1) in Table 2-3 were manufactured at various casting speeds. These billets were heated to 1250°C and expanded to obtain seamless steel pipes. The component compositions of billets No. AO1 to BB1 are the same as those of No. AO to BB shown in Table 1-2. The steel pipes were manufactured under conditions where the expansion was completed at 820°C or higher. For steel pipes with _{an Ac3} point of 950°C or lower, the obtained steel pipes were heated and held at 950°C. For steel pipes with an _Ac3 point exceeding 950°C, the obtained steel pipes were heated and held at 1000°C. After that, they were water-cooled under the conditions described in Table 2-3, and then tempered under the conditions described in Table 2-3. The metallographic structure and mechanical properties were evaluated. The evaluation method is as follows. The tempering temperature was arbitrarily adjusted so that the tensile strength of the material was in the range of 520 MPa to 700 MPa. For the dehydrogenation treatment, the material was held at an ambient temperature (dehydrogenation treatment temperature T) of 50°C for 3 hours, and then allowed to cool naturally.

The fatigue crack propagation characteristics were evaluated by fatigue crack growth tests. Compact tension (CT) test specimens (nearly square specimens with a notch at one end) conforming to ASTM E 647 were taken from each steel material so that the load direction was parallel to the rolling direction. Fatigue tests were performed at a frequency of 1 Hz, with a sinusoidal wave repetition pattern and a stress ratio of R = 0.1. The fatigue crack length was measured using the compliance method with a clip gauge, and the fatigue crack propagation velocity in 5 MPa high-pressure hydrogen gas was determined. The tests were conducted at room temperature (20 ± 10°C). For specimens with a thickness of 10 mm or less, the surface was ground down by 0.5 mm increments to 2 mm, 5 mm, 8 mm, and 9 mm respectively. For other thicknesses, a 10 mm thick specimen was taken from the t/2 (t: thickness) position, and both the front and back surfaces of the crack propagation area were mirror-polished. In this study, the fatigue crack propagation rate (m/cycle) in the stress intensity factor range ΔK = 20 (MPa·m ^1/2 ) was evaluated as a representative value for the stable growth region where Paris's law holds. The results are shown in Tables 2-1, 2-2, and 2-3.

Furthermore, the method for measuring carbides in steel materials is described below. A test specimen was cut from the center of the steel plate thickness, parallel to the thickness direction, and nital etching was performed. Carbides were observed using a scanning electron microscope (SEM). Ten fields of view were randomly selected and observed at an acceleration voltage of 15 kV and a magnification of 20,000x. The average value of the ten fields of view was used as the number of carbides. If there were 20 or fewer carbides with a diameter of 200 nm or more per ¹⁰ μm², it was classified as Y, and if there were more than 20 per 10 ^μm² , it was classified as N, as shown in Tables 2-1, 2-2, and 2-3.
Furthermore, the method for measuring the austenite content of steel materials is as follows.

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 buffed, and then the surface layer was removed by chemical polishing using picric acid etching, and measurements were taken 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.

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 cylindrical in shape with a length of 30 mm in the longitudinal direction of the steel pipe and a diameter of 7Φ, located at the 1/4 position of the thickness of the steel plate and the 1/4 position from the inner surface of the steel pipe. Note that this hydrogen amount is before being subjected to the high-pressure hydrogen fatigue test described later in the aging section, and is the H amount shown in Tables 1-1 and 1-2.

Furthermore, the same tests described above are performed on steel pipes as on other steel materials.

All of the examples of the present invention satisfied the condition that the fatigue crack propagation rate in hydrogen gas was 1.0 × ^10⁻⁶ m/cycle or less.

Furthermore, steel pipes No. 94 and 101, whose casting rates in Table 2-3 were outside the range of the present invention, were used as comparative examples because the formation of coarse inclusions resulted in fatigue crack propagation rates in hydrogen gas being outside the range of the present invention.

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 pipes with the same component composition as billets No. Q and BC shown in Tables 1-1 and 1-2, and AS1 shown in Table 2-3, the cooling process was carried out under predetermined conditions, and after the cooling process (before the tempering process), the steel pipes were reheated under the conditions in Table 3 and the quenching process was performed and their characteristics were evaluated. Steel pipes No. 17A to 17C shown in Table 3 are steel pipes No. 17 shown in Tables 1-1 and 2-1 that have undergone the reheating process. Steel pipes No. 55A to 55C are steel pipes No. 55 shown in Tables 1-2 and 2-2 that have undergone the reheating process, and steel pipes No. 92A to 92B are steel pipes No. 92 shown in Table 2-3 that have undergone the reheating process.
In Example 2, all of the inventive examples satisfied the condition that the crack propagation rate da/dN in hydrogen gas was 1.0 × ^10⁻⁶ m/cycle or less. Among these, the crack propagation characteristics were superior when the reheating and quenching processes were carried out under more favorable conditions.

The following describes examples demonstrating 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 billets No. N and AQ shown in Tables 1-1 and 1-2, and AX1 shown in Table 2-3, the pipes were manufactured under the same conditions as steel pipes No. 14 and 43 shown in Tables 2-1 and 2-2, and steel pipe No. 97 shown in Table 2-3, up to the tempering process, and their characteristics were evaluated when the dehydrogenation treatment conditions were changed. The results are shown in Table 4.
In Example 1, the dehydrogenation treatment of steel pipes No. 14, 43, and 97 was carried out at a dehydrogenation treatment temperature T (ambient temperature) of 50°C for a holding time of 3 hours. However, in this example, for steel pipes No. 14D, 43D, and 97D, the dehydrogenation treatment temperature T (ambient temperature) was set to 50°C, and the holding time tc after the plate thickness center temperature Tc reached 50°C was set to satisfy equation (A). For steel pipes No. 14E, 43E, and 97E, the dehydrogenation treatment temperature T (ambient temperature) was set to 50°C, and the holding time tc was set to satisfy the aforementioned equation (A). However, the holding time tc after the plate thickness center temperature Tc reached 50°C did not satisfy the aforementioned equation (A).
For steel pipes No. 14F and 97F, the dehydrogenation treatment temperature T (ambient temperature) is 50°C, but neither the holding time t at the 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).

The fatigue crack propagation characteristics were evaluated using the fatigue crack growth test described in Example 1.

All of the examples of the present invention satisfied the condition that the crack propagation rate da/dN in hydrogen gas was 1.0 × ^10⁻⁶ m/cycle or less. Among these, the steel pipes in which the dehydrogenation treatment was carried out under more favorable conditions showed superior crack propagation characteristics.

Claims (6)

In mass percent,
C: 0.10-0.45%,
Si: 0.01-2.0%,
Mn: 0.3-2.0%,
Al: 0.01-0.15%,
N: 0.0005-0.008%,
P: 0.015% or less,
S: 0.0015% or less,
O: 0.01% or less,
H: 0.0010% or less,
Cu: 0 to 2.5%,
Ni: 0 to 2.5%,
Cr: 0-2.5%,
Mo: 0-2.0%,
Nb: 0 to 0.5%,
V: 0-0.5%,
Ti: 0 to 0.5%,
W: 0-2.5%,
B: 0 to 0.005%,
Sn: 0-0.3%,
Sb: 0 to 0.3%,
Ca: 0-0.01%,
Mg: 0 to 0.01%,
REM: 0~0.005%
It contains, with the remainder being Fe and unavoidable impurities, and has a component composition.
At a point 1/4 of the wall thickness from the inner surface of the steel pipe, retained austenite is 3% or less by area fraction, and bainite or martensite is 90% or more by area fraction.
The carbides have a diameter of 200 nm or more and are intermetallic compounds containing at least one of cementite, ε-carbide, χ-carbide, or Fe7C3, with 20 or _{fewer carbides} per ¹⁰ μm² .
A method for manufacturing a steel pipe with excellent fatigue properties in hydrogen, wherein the crack propagation rate da/dN in hydrogen at room temperature (20 ± 10°C) and a hydrogen gas pressure of 5 MPa, within a stress intensity factor range of 20 MPa√m, is 1.0 × 10⁻⁶ ^m ·cycle⁻¹ ^or less, comprising a casting step of casting a steel material having the above-mentioned component composition at a casting speed of 1.8 m/min or less,
A heating process that involves heating to 1350°C or below,
A hot rolling process in which the steel material heated in the above heating process is rolled at a rolling completion temperature of 820°C or higher to form a steel pipe shape,
After the steel pipe obtained in the hot rolling process is held at a temperature of Ac ₃ or higher and 1000°C or lower, a cooling process is performed in which the cooling conditions are in group A or group B below.
A method for manufacturing a steel pipe, comprising a tempering step in which the steel pipe obtained in the cooling step is tempered at a temperature of 400°C or higher and an Ac value of ₁ or lower for less than 60 minutes.
Group A:
The average cooling rate from 800°C to 550°C is 15°C/s or more at the 1/4 wall thickness position from the inner surface of the steel pipe, and the average cooling rate from 550°C to 50°C is 15°C/s or less at the 1/4 wall thickness position from the inner surface of the steel pipe, cooling the steel pipe down to 50°C or below.
Group B:
The average cooling rate from 800°C to 300°C is 10°C/s or more at the 1/4 wall thickness position from the inner surface of the steel pipe, and the average cooling rate from 300°C to 50°C is 5°C/s or less at the 1/4 wall thickness position from the inner surface of the steel pipe, cooling the steel pipe down to 50°C or below. A method for manufacturing a steel pipe according to claim 1 , comprising a quenching step before the tempering step, in which the temperature is reheated to an Ac of ₃ or higher and 1000°C or lower, and the cooling conditions are those of group A or group B below.
Group A:
The average cooling rate from 800°C to 550°C is 15°C/s or more at the 1/4 wall thickness position from the inner surface of the steel pipe, and the average cooling rate from 550°C to 50°C is 15°C/s or less at the 1/4 wall thickness position from the inner surface of the steel pipe, cooling the steel pipe down to 50°C or below.
Group B:
The average cooling rate from 800°C to 300°C is 10°C/s or more at the 1/4 wall thickness position from the inner surface of the steel pipe, and the average cooling rate from 300°C to 50°C is 5°C/s or less at the 1/4 wall thickness position from the inner surface of the steel pipe, cooling the steel pipe down to 50°C or below. The method for manufacturing a steel pipe according to claim 1 or 2 , wherein the casting speed is 1.0 m/min or less. In mass percent,
C: 0.10-0.45%,
Si: 0.01-2.0%,
Mn: 0.3-2.0%,
Al: 0.01-0.15%,
N: 0.0005-0.008%,
P: 0.015% or less,
S: 0.0015% or less,
O: 0.01% or less,
H: 0.0010% or less,
Cu: 0 to 2.5%,
Ni: 0 to 2.5%,
Cr: 0-2.5%,
Mo: 0-2.0%,
Nb: 0 to 0.5%,
V: 0-0.5%,
Ti: 0 to 0.5%,
W: 0-2.5%,
B: 0 to 0.005%,
Sn: 0-0.3%,
Sb: 0 to 0.3%,
Ca: 0-0.01%,
Mg: 0 to 0.01%,
REM: 0~0.005%
It contains, with the remainder being Fe and unavoidable impurities, and has a component composition.
At a position 1/4 of the plate thickness from the surface of the steel material, retained austenite accounts for 3% or less by area fraction, and bainite or martensite accounts for 90% or more by area fraction.
The carbides have a diameter of 200 nm or more and are intermetallic compounds containing at least one of cementite, ε-carbide, χ-carbide, or Fe7C3, with 20 or fewer _{carbides per} ¹⁰ μm² .
A method for manufacturing steel with excellent fatigue properties in hydrogen, wherein the crack propagation rate da/dN in hydrogen at room temperature (20 ± 10°C) and hydrogen gas pressure of 5 MPa, within a stress intensity factor range of 20 MPa√m, is 1.0 × 10⁻⁶ ^m ·cycle⁻¹ ^or less, comprising a casting step of casting a steel material having the above-mentioned component composition at a casting speed of 1.8 m/min or less,
A heating process that involves heating to 1350°C or below,
A hot rolling process is performed in which the steel material heated in the aforementioned heating process is rolled at a rolling completion temperature of 820°C or higher.
The steel material obtained in the hot rolling process is held at a temperature of Ac ₃ or higher and 1000°C or lower, followed by a cooling process where the cooling conditions are in group A or group B below.
A method for manufacturing steel, comprising: a tempering step in which the steel obtained in the cooling step is tempered at a temperature of 400°C or higher and an Ac of ₁ point or lower for less than 60 minutes.
Group A:
The average cooling rate from 800°C to 550°C is 15°C/s or more at a point 1/4 of the thickness from the surface of the steel material, and the average cooling rate from 550°C to 50°C is 15°C/s or less at a point 1/4 of the thickness from the surface of the steel material, cooling the steel material down to 50°C or below.
Group B:
The average cooling rate from 800°C to 300°C is 10°C/s or more at a point 1/4 of the thickness from the surface of the steel material, and the average cooling rate from 300°C to 50°C is 5°C/s or less at a point 1/4 of the thickness from the surface of the steel material, thereby cooling the steel material to 50°C or below. A method for manufacturing steel according to claim 4 , comprising a quenching step before the tempering step, in which the steel is reheated to an Ac of ₃ or higher and 1000°C or lower, and the cooling conditions are those of group A or group B below.
Group A:
The average cooling rate from 800°C to 550°C is 15°C/s or more at a point 1/4 of the thickness from the surface of the steel material, and the average cooling rate from 550°C to 50°C is 15°C/s or less at a point 1/4 of the thickness from the surface of the steel material, cooling the steel material down to 50°C or below.
Group B:
The average cooling rate from 800°C to 300°C is 10°C/s or more at a point 1/4 of the thickness from the surface of the steel material, and the average cooling rate from 300°C to 50°C is 5°C/s or less at a point 1/4 of the thickness from the surface of the steel material, thereby cooling the steel material to 50°C or below. The method for manufacturing steel according to claim 4 or 5 , wherein the casting speed is 1.0 m/min or less.