WO2025177987A1 - High-strength steel sheet, steel pipe, method for producing high-strength steel sheet, and method for producing steel pipe - Google Patents

Abstract

The present invention provides a high-strength steel sheet serving as a material for steel pipes satisfying all of the following requirements: large diameter, excellent low-temperature toughness, and possessing strength characteristics required for high-strength steel pipes for vapor transport. The present invention also provides a method for producing the same. In addition, the present invention provides: a steel pipe composed of the high-strength steel sheet; and a method for producing the same.　A high-strength steel sheet according to the present invention has a specific component composition and a specific microstructure; the toughness of the steel sheet is 85% or greater in terms of the ductile fracture rate obtained by DWTT at -40°C; and the yield strength before and after aging performed under the condition of LMP = 15700 is 555 MPa or greater, LMP being a Larson Miller Parameter, defined by formula (2). Formula (2): LMP = (T + 273) × (20 + log(t)), where T is the heat treatment temperature (°C) and t is the heat treatment time (time)

Description

High-strength steel plate, steel pipe, method for manufacturing high-strength steel plate, and method for manufacturing steel pipe

The present invention relates to a high-strength steel plate having excellent low-temperature toughness and a yield strength of 555 MPa or more before and after long-term aging in the medium temperature range, a method for manufacturing the same, and a steel pipe made from the high-strength steel plate and a method for manufacturing the same. In particular, the high-strength steel plate of the present invention is suitable for use as a material for high-strength steel pipes for steam piping.

Methods for recovering oil sands from oil reservoirs in Canada and elsewhere include open-cut mining and the steam injection method, in which high-temperature, high-pressure steam is pumped into the reservoir through steel pipes. Because open-cut mining is not feasible in many areas, the steam injection method is used in many regions. In the steam injection method, the steam pumped into the reservoir is in the 300-400°C temperature range (hereinafter referred to as the medium-temperature range). In the steam injection method, medium-temperature steam is pumped into the reservoir at high pressure. As mentioned above, steel pipes are used to pump this steam. In recent years, with the rise in energy demand, there has been a demand for larger diameter and higher strength steel pipes to improve heavy oil recovery rates and reduce construction costs. Furthermore, when laying the steel pipes, construction work and hydraulic pressure testing are typically carried out from early spring onward when temperatures are high. However, due to increased demand, there is also a demand for pipes with excellent low-temperature toughness to enable construction work during the cooler winter months.

Patent Documents 1 and 2 are examples of prior art steel pipes for transporting steam that can be used in the steam injection method. These patent documents disclose seamless steel pipes equivalent to API X80 grade, with a maximum outer diameter of 16 inches. It is difficult to further increase the diameter of seamless steel pipes. Furthermore, in order to achieve strength equal to or greater than API X80 grade, the addition of large amounts of alloying elements is required.

Patent Documents 3 and 4 disclose manufacturing technologies for high-strength steel pipes that can be made larger in diameter and are manufactured by welding. More specifically, Patent Documents 3, 4, and 5 relate to manufacturing technologies for high-strength steel pipes that are produced by TMCP (thermo-mechanical control process) and have strengths of API X80 grade or higher.

Japanese Patent Application Laid-Open No. 2000-290728 Patent No. 4821939 Patent No. 5055736 Patent No. 4741528 Patent No. 6241569

Steel pipes manufactured using the steel pipe manufacturing method disclosed in Patent Document 3 satisfy API X80 grade high-temperature properties in the medium temperature range. However, Patent Document 3 does not disclose the strength properties of the steel pipes when used for long periods of time in the medium temperature range.

Patent Document 4 describes the creep properties of steel pipes after long-term aging, and while the developed steel shows improvements in breaking strength, the safety criteria are not clearly defined. Furthermore, there are also instances where breaking strength values are below 80% (= 440 MPa) of the lower limit of the standard for yield stress in API X80 grade, making it difficult to say that the strength is sufficient.

The steel pipe disclosed in Patent Document 5 satisfies the API X80 grade in terms of high-temperature properties in the medium temperature range, but has a small LMP, and no consideration is given to the strength properties when the steel pipe is used for long periods of time, nor to low-temperature toughness.

As such, conventional technology is unable to obtain high-strength steel pipe for steam piping that meets all of the following requirements: large diameter, strength characteristics required for high-strength steel pipe for steam transportation, and excellent low-temperature toughness.

The present invention has been made to solve the above problems, and aims to provide a high-strength steel plate that can be used as a material for steel pipes that meet all of the requirements for large diameter, excellent low-temperature toughness, and strength characteristics required for high-strength steel pipes for steam transportation, as well as a manufacturing method for the same. The present invention also aims to provide a steel pipe made from the above-mentioned high-strength steel plate and a manufacturing method for the same.

Furthermore, "excellent low-temperature toughness" means that the ductile fracture surface area ratio (DWTTSA-40°C) obtained by DWTT (test temperature: -40°C) in accordance with API 5L is 85% or more, and the fracture surface transition temperature is -40°C or less. The test temperature for DWTT is set at -40°C to take into account the reduction in toughness due to work hardening during pipe manufacturing.

The inventors conducted extensive research into the properties of high-strength steel plate for large-diameter steel pipes in the medium temperature range. As a result, they discovered that by appropriately selecting the chemical composition and manufacturing conditions, it is possible to obtain high-strength steel plate that can be used to manufacture high-strength steel pipes that, despite their large diameter, have the strength characteristics and low-temperature toughness required for high-strength steel pipes used to transport steam.

The present invention was made based on the above findings and further studies, and is therefore comprised of the following:
[1] The component composition is, in mass%,
C: 0.04-0.09%,
Si: 0.03-0.25%,
Mn: 1.5-2.5%,
P: 0.020% or less,
S: 0.002% or less,
Mo: 0.10-0.50%,
Nb: 0.010-0.055%,
Ti: 0.005-0.020%,
Ca: 0.0040% or less,
Al: 0.01-0.04%,
N: 0.006% or less, the balance being Fe and unavoidable impurities;
X (%) represented by formula (1) is 0.65% or more,
the microstructure has bainite in an area ratio of 80% or more at the center of the plate thickness, the average grain size of the bainite is 25 μm or less, and the minimum grain size of the top 20% of the bainite grains with the largest grain size is 60 μm or less, the toughness of the steel plate is such that the ductile fracture area ratio obtained by DWTT at -40°C is 85% or more,
A high-strength steel plate having a yield strength of 555 MPa or more before and after aging under the condition of Larson Miller Parameter LMP = 15700, as defined by the following formula (2).
X (%) = 0.35Cr + 0.9Mo + 12.5Nb + 8V... (1)
The element symbols in formula (1) represent the content (mass%) of each element. Elements that are not contained are substituted with 0.
LMP=(T+273)×(20+log(t))...(2)
T: Heat treatment temperature (°C)
t: heat treatment time (hours)
[2] The component composition, in mass%, further comprises:
Cr: 0.50% or less,
V: 0.070% or less,
Cu: 0.50% or less,
The high-strength steel plate according to [1], containing one or more of: Ni: 0.50% or less.
[3] Further, the sheet has an area ratio of 10% or less of ferrite and an area ratio of 10% or less of island martensite at the center of the sheet thickness,
the total number of Ti-based precipitates having a diameter of 100 nm or less, Nb-based precipitates having a diameter of 100 nm or less, V-based precipitates having a diameter of 100 nm or less, Mo-based precipitates having a diameter of 100 nm or less, Cr-based precipitates having a diameter of 100 nm or less, Al-based precipitates having a diameter of 100 nm or less, and composite precipitates having a diameter of 100 nm or less containing two or more elements selected from Ti, Nb, V, Mo, Cr, and Al is 50,000 to 1,000,000 ^per 1 mm2 of the test area, at the center of the sheet thickness;
3. The high-strength steel plate according to claim 1, wherein a difference in yield strength before and after aging, obtained by subtracting the yield strength after aging from the yield strength before aging, is 50 MPa or less.
[4] A steel pipe using the high-strength steel plate according to any one of [1] to [3].
[5] A method for producing a high-strength steel plate according to any one of [1] to [3], comprising: a heating step of heating a steel material to 1000 to 1200 ° C.;
a hot rolling process in which the steel material heated in the heating process is hot rolled under the conditions of one or more passes of rolling at 950°C or higher with a rolling reduction of 10% or more per pass and one or more passes of rolling at 900°C or lower with a rolling reduction of 15% or more per pass, a cumulative rolling reduction of 50% or more at 900°C or lower, and a rolling end temperature of _Ar3 temperature or higher and 850°C or lower;
and an accelerated cooling step of accelerated cooling the hot-rolled steel sheet obtained in the hot rolling step under conditions of a cooling start temperature of Ar ₃ temperature or higher, an average cooling rate of 5°C/s or higher, and a cooling stop temperature of 300 to 550°C.
[6] A cold forming step of cold forming the high-strength steel plate according to any one of [1] to [3] into a tubular shape;
a welding process for welding the butt joints at the butt joints of the ends of the steel plates formed into a tubular shape in the cold forming process;
A method for manufacturing a steel pipe having the above structure.

The present invention makes it possible to obtain high-strength steel pipes that have the strength and toughness characteristics required for high-strength steel pipes used to transport steam, despite their large diameter.

The following describes an embodiment of the present invention. Note that the present invention is not limited to the following embodiment.

<High strength steel plate>
The high-strength steel plate (also simply referred to as steel plate) of the present invention contains, in mass %, 0.04 to 0.09% C, 0.03 to 0.25% Si, 1.5 to 2.5% Mn, 0.020% or less P, 0.002% or less S, 0.10 to 0.50% Mo, 0.010 to 0.055% Nb, 0.005 to 0.020% Ti, 0.0040% or less Ca, 0.01 to 0.04% Al, and 0.006% or less N. In the following description, "%" representing the content of a component means "% by mass."

C: 0.04-0.09%
C is an element necessary for ensuring the strength of steel through solid solution strengthening and precipitation strengthening. In particular, an increase in the amount of solute C and the formation of precipitates contribute to ensuring strength in the mid-temperature range. In order to ensure a predetermined strength at room temperature and the mid-temperature range, the present invention specifies a C content of 0.04% or more. The C content is preferably 0.05% or more. Furthermore, a C content exceeding 0.09% leads to deterioration of toughness and weldability. For this reason, the upper limit of the C content is specified as 0.09%. Furthermore, the C content is preferably 0.08% or less, and more preferably 0.07% or less.

Si: 0.03-0.25%
Si is added for deoxidation. If the Si content is less than 0.03%, a sufficient deoxidation effect cannot be obtained, and the toughness of the steel plate decreases. Therefore, the Si content is set to 0.03% or more. The Si content is preferably 0.04% or more, more preferably 0.05% or more, even more preferably 0.07% or more, and most preferably 0.10% or more. On the other hand, if the Si content exceeds 0.25%, the toughness will deteriorate. Therefore, the Si content is set to 0.25% or less. Furthermore, the Si content is preferably 0.23% or less, more preferably 0.20% or less, even more preferably 0.19% or less, and most preferably 0.18% or less.

Mn: 1.5-2.5%
Mn is an element effective in improving the strength and toughness of steel. If the Mn content is less than 1.5%, the effect is small. Therefore, the Mn content is set to 1.5% or more. The Mn content is preferably set to 1.55% or more, more preferably 1.6% or more, even more preferably 1.65% or more, and most preferably 1.7% or more. On the other hand, if the Mn content exceeds 2.5%, toughness and weldability significantly deteriorate. Therefore, the Mn content is set to 2.5% or less. Furthermore, the Mn content is preferably set to 2.4% or less, more preferably 2.3% or less, even more preferably 2.2% or less, and most preferably 2.1% or less.

P: 0.020% or less P is an impurity element that deteriorates toughness. For this reason, it is desirable to reduce the P content as much as possible. However, excessive reduction of the P content increases manufacturing costs. Therefore, the P content is set to 0.020% or less under the condition that the deterioration of toughness falls within an acceptable range. The P content is preferably set to 0.018% or less, more preferably 0.015% or less, even more preferably 0.013% or less, and most preferably 0.010% or less. The lower limit of the P content is not particularly limited, and may be 0%, but is more preferably 0.003% or more.

S: 0.002% or less S is an impurity element that deteriorates toughness. For this reason, it is desirable to reduce the S content as much as possible. Furthermore, even if Ca is added to control the shape of MnS to CaS-based inclusions, in the case of X80 grade high-strength materials, finely dispersed CaS-based inclusions can also be a factor in reducing toughness. For this reason, the S content is set to 0.002% or less. The S content is preferably set to 0.0018% or less, more preferably 0.0015% or less, even more preferably 0.0013% or less, and most preferably 0.0010% or less. The lower limit of the S content is not particularly limited, and may be 0%, but is preferably 0.0003% or more.

Mo: 0.10~0.50%
Mo solid-solution strengthens steel and improves its hardenability. Solid-solution strengthening and improved hardenability increase strength, particularly in the intermediate temperature range, due to increased tempering softening resistance. If the Mo content is less than 0.10%, the effect is small and sufficient strength cannot be obtained. Therefore, the Mo content is set to 0.10% or more. The Mo content is preferably set to 0.13% or more, more preferably 0.15% or more, even more preferably 0.18% or more, and most preferably 0.20% or more. On the other hand, if the Mo content exceeds 0.50%, the effect saturates and toughness and weldability deteriorate. Therefore, the Mo content is set to 0.50% or less. Furthermore, the Mo content is preferably set to 0.45% or less, more preferably 0.40% or less, even more preferably 0.35% or less, and most preferably 0.30% or less.

Nb: 0.010-0.055%
Nb suppresses grain growth during slab heating and rolling. This refines the microstructure and provides the steel with sufficient strength and toughness. Nb is also an essential component for forming carbides to ensure strength in the intermediate temperature range. This effect is most pronounced at an Nb content of 0.010% or more. Therefore, the Nb content is set to 0.010% or more. The Nb content is preferably set to 0.013% or more, more preferably 0.015% or more, even more preferably 0.018% or more, and most preferably 0.020% or more. On the other hand, if the Nb content exceeds 0.055%, not only does the effect saturate, but toughness and weldability also deteriorate. Therefore, the Nb content is set to 0.055% or less. The Nb content is preferably set to 0.053% or less, more preferably 0.050% or less, even more preferably 0.048% or less, and most preferably 0.045% or less.

Ti: 0.005-0.020%
Ti forms TiN to suppress grain growth during slab heating and in weld heat-affected zones, thereby improving toughness by refining the microstructure. To achieve this effect, the Ti content must be 0.005% or more. The Ti content is preferably 0.006% or more, more preferably 0.007% or more, even more preferably 0.008% or more, and most preferably 0.009% or more. On the other hand, if the Ti content exceeds 0.020%, toughness deteriorates. Therefore, the Ti content is set to 0.020% or less. The Ti content is preferably 0.019% or less, more preferably 0.018% or less, even more preferably 0.017% or less, and most preferably 0.016% or less.

Ca: 0.0040% or less Ca functions to fix S in steel and improve the toughness of the steel plate. To achieve this effect, the Ca content is preferably 0.0010% or more, and more preferably 0.0015% or more. However, if the Ca content exceeds 0.0040%, inclusions in the steel increase, which may deteriorate the toughness. Therefore, the Ca content is set to 0.0040% or less. The Ca content is preferably 0.0038% or less, more preferably 0.0035% or less, even more preferably 0.0032% or less, and most preferably 0.0030% or less.

Al: 0.01-0.04%
Al is added as a deoxidizer. If the Al content is less than 0.01%, a sufficient deoxidizing effect cannot be obtained, and the toughness of the steel plate decreases. The Al content is set to 0.01% or more, preferably 0.013% or more, more preferably 0.015% or more, even more preferably 0.018% or more, and most preferably 0.02% or more. On the other hand, if the Al content exceeds 0.04%, the toughness deteriorates. Therefore, the Al content is set to 0.04% or less. The Al content is set to 0.038% or less, more preferably 0.035% or less, even more preferably 0.032% or less, and most preferably 0.03% or less.

N: 0.006% or less N forms TiN together with Ti and is finely dispersed in the high-temperature region of the weld heat-affected zone, which reaches 1350°C or higher. This fine dispersion refines the prior austenite grains in the weld heat-affected zone, significantly contributing to improving the toughness of the weld heat-affected zone. To achieve this effect, the N content is preferably 0.002% or more, and more preferably 0.0025% or more. On the other hand, if the N content exceeds 0.006%, the base material toughness will deteriorate due to coarsening of precipitates and an increase in solute N, and the toughness of the weld metal in steel pipes will deteriorate. Therefore, the N content is set to 0.006% or less. The N content is preferably 0.0058% or less, more preferably 0.0055% or less, even more preferably 0.0053% or less, and most preferably 0.005% or less.

X (X(%) = 0.35Cr + 0.9Mo + 12.5Nb + 8V (1)): 0.65% or more In the present invention, the Cr content, Mo content, Nb content, and V content are adjusted so that X(%), represented by the following formula (1), is 0.65% or more. X(%) is an important factor for obtaining a steel having excellent strength after long-term aging in the intermediate temperature range. Therefore, X(%) needs to be 0.65% or more to suppress dislocation recovery during long-term aging. X(%) is preferably 0.66% or more, more preferably 0.67% or more, even more preferably 0.68% or more, and most preferably 0.69% or more. In order to inexpensively produce the high-strength steel sheet of the present invention, X(%) is preferably 1.20% or less, more preferably 1.19% or less, even more preferably 1.18% or less, and most preferably 1.17% or less. In addition, there are cases in which Cr and V are not included in the present invention, in which case 0 should be substituted for "Cr" and "V" in formula (1).
X (%) = 0.35Cr + 0.9Mo + 12.5Nb + 8V... (1)
The element symbols in formula (1) represent the content (mass%) of each element. Elements that are not contained are substituted with 0.

Furthermore, the high-strength steel sheet of the present invention may contain one or more of Cr, V, Cu, and Ni to further improve its properties.

Cr: 0.50% or less Cr is one of the elements effective in increasing temper softening resistance and high-temperature strength. On the other hand, a Cr content exceeding 0.50% adversely affects weldability. Therefore, when Cr is contained, the Cr content is set to 0.50% or less. The Cr content is preferably set to 0.48% or less, more preferably 0.45% or less, even more preferably 0.43% or less, and most preferably 0.40% or less. Note that, as long as the temper softening resistance and high-temperature strength can be increased by, for example, ensuring that X is within the desired range, the lower limit of the Cr content is not particularly limited and may be 0%. To achieve the above effect, the Cr content is preferably 0.05% or more, more preferably 0.08% or more, and even more preferably 0.10% or more.

V: 0.070% or less A small amount of V refines grains and contributes to increased strength. It also increases temper softening resistance and is one of the elements effective in increasing strength in the medium temperature range. On the other hand, if the V content exceeds 0.070%, the toughness of the weld heat-affected zone deteriorates. Therefore, when V is contained, the V content is set to 0.070% or less. The V content is preferably 0.060% or less, more preferably 0.050% or less, even more preferably 0.040% or less, and most preferably 0.030% or less. Note that, as long as high-temperature strength can be increased by, for example, ensuring that X is within the desired range, the lower limit of the V content is not particularly limited and may be 0%. To achieve the above effect, the V content is preferably 0.005% or more, more preferably 0.010% or more.

Cu: 0.50% or less Cu is one of the elements effective in improving toughness and increasing strength. To obtain these effects, the Cu content is preferably 0.05% or more, and more preferably 0.10% or more. On the other hand, if the Cu content exceeds 0.50%, weldability decreases. Therefore, when Cu is contained, the Cu content is set to 0.50% or less. The Cu content is preferably 0.48% or less, more preferably 0.45% or less, even more preferably 0.43% or less, and most preferably 0.40% or less.

Ni: 0.50% or less Ni is one of the elements effective in improving toughness and increasing strength. To obtain these effects, the Ni content is preferably 0.05% or more, and more preferably 0.10% or more. If the Ni content exceeds 0.50%, the effect saturates and production costs increase. Therefore, when Ni is contained, the Ni content is set to 0.50% or less. The Ni content is preferably 0.48% or less, more preferably 0.45% or less, even more preferably 0.43% or less, and most preferably 0.40% or less.

In the present invention, it is preferable to adjust the Cu content, Ni content, Cr content, and Mo content so that Cu + Ni + Cr + Mo (the element symbols represent the content (mass%) of each element) is 1.00% or less. In order to achieve both the strength-increasing effect of these elements and manufacturing costs, Cu + Ni + Cr + Mo is preferably 1.00% or less, more preferably 0.98% or less, even more preferably 0.95% or less, and most preferably 0.90% or less. There is no particular lower limit, but Cu + Ni + Cr + Mo is preferably 0.10% or more, more preferably 0.15% or more, even more preferably 0.20% or more, and most preferably 0.25% or more.

The remainder of the components is Fe and unavoidable impurities. These impurities are unavoidably mixed in from raw materials, the manufacturing process, or manufacturing equipment, and are permitted to be present to the extent that they do not impair the objectives of the present invention. Raw materials include iron ore, reduced iron, and scrap. Examples of unavoidable impurities include Pb, Zn, Sn, As, B, Sb, Bi, Co, H, O, and REM.

Next, we will explain the structure of the high-strength steel plate of the present invention.

Bainite is an important structure for achieving both strength and low-temperature toughness. Bainite effectively contributes to improving the strength of steel plates by strengthening the transformation structure. Microstructural uniformity is necessary for increasing the initial dislocation density and for improving the strength of high-strength steels, particularly in the mid-temperature range. Therefore, the structure of the high-strength steel of the present invention must be predominantly bainite. Specifically, the area fraction of bainite relative to the entire steel structure at the thickness center must be 80% or more. The area fraction of bainite is preferably 83% or more, more preferably 85% or more, even more preferably 88% or more, and most preferably 90% or more. While there are no particular limitations on the upper limit of the bainite fraction, from the viewpoint of improving deformability, the area fraction of bainite is preferably 98% or less, more preferably 95% or less.

Ferrite: area ratio of 10% or less (optimal condition)
When ferrite is processed by rolling to become processed ferrite, the number of mobile dislocations increases, which can increase strength. On the other hand, the dislocation density of the ferrite phase is significantly reduced by heat treatment in the medium temperature range, which softens the ferrite phase and may reduce the strength of the steel sheet after heat treatment. Therefore, it is preferable that the area fraction of ferrite at the center of the sheet thickness is 10% or less. The area fraction of ferrite is more preferably 9% or less, even more preferably 8% or less, and most preferably 7% or less. There is no particular limit on the lower limit of the area fraction of ferrite, but from the viewpoint of improving deformability, the area fraction of ferrite is preferably 1% or more.

Island martensite (MA): area ratio of 10% or less (optimal condition)
Island martensite (MA: Martensite-Austenite constituent) is a very hard phase and may act as a fracture initiation point, thereby reducing the low-temperature toughness of the steel plate. Therefore, the area fraction of island martensite (MA) at the center of the plate thickness is preferably 10% or less. The area fraction of island martensite is more preferably 8% or less, even more preferably 5% or less, and most preferably 3% or less. There is no particular restriction on the lower limit of the area fraction of island martensite, but the area fraction is preferably 1% or more, and more preferably 2% or more.
As described above, the steel structure of the steel plate serving as the base material must basically be composed of the above-mentioned bainite, but examples of the remaining structure other than bainite, ferrite, and island martensite (MA) include pearlite, martensite, cementite, and retained austenite.

The average grain size of bainite is 25 μm or less, and the minimum grain size of the top 20% of bainite grains with the largest grain size is 60 μm or less. Because bainite grain boundaries provide resistance to brittle crack propagation, grain refinement contributes to improving low-temperature toughness. Therefore, the average grain size of all bainite is 25 μm or less. The average grain size of all bainite is preferably 23 μm or less, more preferably 20 μm or less, even more preferably 19 μm or less, and most preferably 18 μm or less. There is no particular lower limit, but 5 μm or more is preferred, and 6 μm or more is more preferred.

Furthermore, while refining the average grain size improves low-temperature toughness, there is a limit to how much the average grain size can be refined when cooling is initiated near the _Ar3 temperature. In the present invention, it is essential to suppress the formation of even larger grains. Coarse bainite is likely to become the origin of fracture, and if the top 20% of bainite grains have large grain sizes, low-temperature toughness deteriorates. In particular, if the minimum grain size of the top 20% of grains with the largest grain size exceeds 60 μm, they are likely to become the origin of fracture. Therefore, the minimum grain size of the top 20% of grains with the largest grain size at the center of the plate thickness must be 60 μm or less, preferably 55 μm or less, more preferably 50 μm or less, even more preferably 45 μm or less, and most preferably 40 μm or less. The lower limit is not particularly limited, but is preferably 10 μm or more, and more preferably 15 μm or more.

Here, the average grain size of bainite is determined as follows.
That is, after mirror polishing the L cross section of the steel sheet (a cross section parallel to the rolling direction and parallel to the normal direction of the rolling surface), the crystal orientation of a randomly selected 1 mm x 1 ^mm2 region at the center of the sheet thickness was measured by electron backscatter diffraction (EBSD), and regions where the angle difference between adjacent pixels was 15° or more were determined as grain boundaries by image analysis. The EBSD measurement conditions were an acceleration voltage of 17 kV and a measurement pitch of 0.8 μm.
The average crystal grain size darea (μm) was calculated from the area ai (μm ² ) occupied by each crystal grain and the circle-equivalent diameter di (μm) of each crystal grain using the following formula:
darea (μm) = Σ(ai・di)/Σai
Furthermore, the minimum grain size of the top 20% of bainite grains by grain size represents the grain size of the smallest grain when the circle equivalent diameters of the grains are arranged in descending order of grain size and 20% of the total number of grains are selected from the largest.

The total number of Ti-based precipitates with a diameter of 100 nm or less, Nb-based precipitates with a diameter of 100 nm or less, V-based precipitates with a diameter of 100 nm or less, Mo-based precipitates with a diameter of 100 nm or less, Cr-based precipitates with a diameter of 100 nm or less, Al-based precipitates with a diameter of 100 nm or less, and composite precipitates with a diameter of 100 nm or less consisting of two or more elements selected from Ti, Nb, V, Mo, Cr, and Al is 50,000 to 1,000,000 per 1 ^mm2 of the test area (optimal conditions).
The reasons for limiting the sizes and total number of Ti-based precipitates, Nb-based precipitates, V-based precipitates, Mo-based precipitates, Cr-based precipitates, Al-based precipitates, and composite precipitates containing two or more elements of Ti, Nb, V, Mo, Cr, and Al (for example, one or more of Ti-Nb-based precipitates, Ti-V-based precipitates, Cr-Mo-based precipitates, Cr-Mo-Nb-based precipitates, Al-Ti-based precipitates, Al-Nb-based precipitates, and Al-Ti-Nb-based precipitates) in the steel of the present invention will be explained below.
Even if precipitates with a diameter of 100 nm or less (fine precipitates) are formed, if the number of such precipitates is less than 50,000 per ^mm2 , the austenite grain growth inhibitory effect is weak, and strength degradation in the intermediate temperature range may not be suppressed. Therefore, it is preferable that the number of fine precipitates with a diameter of 100 nm or less in the steel at the center of the plate thickness is 50,000 or more ^per mm2. More preferably, it is 80,000 or more, even more preferably, it is 100,000 or more, and most preferably, it is 130,000 or more. On the other hand, if the number of precipitates with a diameter of 100 nm or less exceeds 1,000,000 per ^mm2 , the fine precipitates may aggregate and coarsen, thereby weakening the austenite grain growth inhibitory effect and increasing strength degradation in the intermediate temperature range. Therefore, it is preferable that the number of fine precipitates with a diameter of 100 nm or less in the steel be 1,000,000 or less ^per mm2. More preferably, it is 950,000 or less, even more preferably, 900,000 or less, and most preferably, 850,000 or less.

Here, we will explain how to measure the density and size (diameter) of precipitates. The corroded surface at any location on the head cross section is observed using a scanning electron microscope (SEM), or an extracted replica sample or thin film sample is prepared and observed using a transmission electron microscope (TEM). The number of precipitates with a size of 100 nm or less is measured in an area of at least 100 ^μm2 . For example, when observing at a magnification of 100,000 times with one field of view being 2000 nm x 2000 nm, the observation area per field is 4 ^μm2 , so 25 fields of view are observed randomly. This measurement result is converted to the number per unit area. If the number of precipitates with a size of 100 nm or less is 5 in 25 fields of view (100 ^μm2 ), the density of precipitates can be converted to 50,000 precipitates per ^mm2 . The density of the precipitates is called the number per ^mm2 of the observation area. The diameter of the precipitate is the average value of the major axis (long side) and minor axis (short side).

Next, the strength properties of the high-strength steel sheet of the present invention will be described. The high-strength steel sheet of the present invention has a yield strength of 555 MPa or more before and after aging under the condition of LMP=15700, which is a Larson Miller parameter.
Aging treatment with an LMP of 15700 refers to aging treatment performed under conditions of a heat treatment temperature and a heat treatment time such that the LMP, expressed by the following formula (2), becomes 15700. An LMP of 15700 corresponds to a condition of heat treatment at 350°C, which is a medium temperature range, for 20 years. The condition of Larson Miller Parameter (LMP) of 15700 means that the Larson Miller Parameter (LMP) is equal to or greater than 15650 and less than 15750.
LMP=(T+273)×(20+log(t))...(2)
T: Heat treatment temperature (°C)
t: heat treatment time (hours)
The high-strength steel sheet of the present invention has a yield strength of 555 MPa or more before and after the aging treatment. A yield strength of 555 MPa or more has the effect of enabling stable operation as a steel pipe for steam piping. Here, the yield strength means the yield strength measured in a high-temperature tensile test at 350°C. The yield strength before and after the aging treatment is preferably 560 MPa or more, more preferably 565 MPa or more, even more preferably 570 MPa or more, and most preferably 575 MPa or more. Although there is no particular upper limit, the yield strength is preferably 840 MPa or less.
As described in the Examples, in the present invention, the test specimens taken from both the steel plate and the steel pipe have a yield strength of 555 MPa or more both before and after the aging treatment.

The difference in yield strength before and after aging, calculated by subtracting the yield strength after aging from the yield strength before aging, is 50 MPa or less (suitable requirement).
In the present invention, it is preferable that the difference between the yield strength at 350°C measured after aging under the condition of Larson Miller Parameter (LMP) = 15700 and the yield strength at 350°C measured before the aging satisfies the relationship of 50 MPa or less. The yield strength before and after aging is an index for evaluating the decrease in yield strength when held for a long time in the intermediate temperature range. If this difference is 50 MPa or less, the decrease in yield strength after long-term holding in the intermediate temperature range is within a range that is practically acceptable. If this difference is greater than 50 MPa, a significant decrease in yield strength occurs after long-term holding in the intermediate temperature range. Therefore, it is preferable that the difference in yield strength before and after aging be 50 MPa or less. The difference in yield strength before and after aging is more preferably 45 MPa or less, even more preferably 40 MPa or less, and most preferably 35 MPa or less. The lower limit is not particularly limited and may be a negative value, such as -100 MPa or more.
The aging treatment conditions under the above conditions include, for example, heat treatment at 400° C. for 2335 hours.

The toughness of the steel plate is such that the ductile fracture surface area ratio (DWTT) obtained by DWTT at -40°C is 85% or more. The toughness of the steel plate of the present invention is such that the ductile fracture surface area ratio (DWTTSA-40°C) obtained by DWTT (test temperature: -40°C) in accordance with API 5L is 85% or more. If the ductile fracture surface area ratio is less than 85%, the steel plate is prone to brittle fracture at low temperatures, making it difficult to install the steel plate throughout the year, including in winter when the temperature drops below 0°C, or to use it in areas with very low ambient temperatures. Therefore, the ductile fracture surface area ratio must be 85% or more. Furthermore, a ductile fracture surface area ratio of 85% or more obtained by DWTT at -40°C means that the fracture transition temperature is -40°C or less. The ductile fracture surface area ratio is preferably 86% or more, more preferably 87% or more, even more preferably 88% or more, and most preferably 89% or more. The reason why the test temperature in the DWTT is set to −40° C. is to take into account the decrease in toughness due to work hardening during pipe making. The upper limit of the ductile fracture area ratio obtained by the DWTT at −40° C. is not particularly limited, and may be 100% or less.

<Steel pipe>
The steel pipe of the present invention is manufactured using the high-strength steel plate of the present invention, and therefore has the strength characteristics and low-temperature toughness required for high-strength welded steel pipe for steam transportation, even when it has a large diameter.

Large diameter means that the outer diameter (diameter) of the steel pipe is 400 mm or more. The outer diameter of the steel pipe is preferably 500 mm or more, more preferably 600 mm or more, and even more preferably 700 mm or more. There are no particular restrictions on the maximum outer diameter, but it may be 1500 mm or less, and more preferably 1400 mm or less. According to the present invention, it is possible to increase the diameter while maintaining the strength characteristics required of high-strength welded steel pipes for steam transportation.

Furthermore, the thickness of the steel pipe is not particularly limited, but in the case of steam transport, it is 12 to 30 mm. That is, the thickness of the steel pipe is preferably 12 mm or more, more preferably 13 mm or more, even more preferably 14 mm or more, and most preferably 15 mm or more. The thickness of the steel pipe is preferably 30 mm or less, more preferably 29 mm or less, even more preferably 28 mm or less, and most preferably 27 mm or less.

The strength characteristics required for high-strength welded steel pipe for steam transportation are, like the above-mentioned high-strength steel, a yield strength at 350°C of 555 MPa or more before and after aging under the condition of Larson Miller Parameter (LMP) = 15700. The yield strength is preferably 560 MPa or more, more preferably 565 MPa or more, even more preferably 570 MPa or more, and most preferably 575 MPa or more.
Although there is no particular upper limit, the yield strength at 350°C before and after the aging is preferably 840 MPa or less, and more preferably 830 MPa or less.

<Method of manufacturing high-strength steel plate>
Next, a method for manufacturing a high-strength steel plate according to the present invention will be described. The method for manufacturing a high-strength steel plate according to the present invention includes a heating step, a hot rolling step, and an accelerated cooling step. Each step will be described below. In the following description, unless otherwise specified, temperatures refer to the average temperature in the thickness direction. The average temperature in the thickness direction is determined by simulation calculations or the like using parameters such as the plate thickness, cooling conditions, and heat transfer coefficient from the surface temperature of the slab or steel plate. For example, the average temperature in the thickness direction can be determined by calculating the temperature distribution in the thickness direction using a finite difference method. The average temperature in the thickness direction is the average value of each temperature in the thickness direction from the surface of the steel plate to the surface of the opposite steel plate. The cooling rate is the average cooling rate obtained by dividing the temperature difference required for cooling to the cooling stop (end) temperature after the end of hot rolling by the time required for that cooling.

Heating Step In the method for producing a high-strength steel plate of the present invention, the heating step is a step of heating a steel material to 1000 to 1200°C. Here, the steel material is, for example, a slab obtained by casting molten steel. Since the chemical composition of the steel material will be the chemical composition of the high-strength steel plate, the chemical composition of the high-strength steel plate can be adjusted at the stage of adjusting the chemical composition of the molten steel. Note that there are no particular limitations on the steelmaking method for the steel material.

During the hot rolling process described below, in order to sufficiently promote austenitization and carbide solid solution and obtain sufficient strength at room temperature and in the mid-temperature range, the heating temperature of the steel material must be 1000°C or higher. The heating temperature is preferably 1010°C or higher, more preferably 1020°C or higher, even more preferably 1030°C or higher, and most preferably 1040°C or higher. On the other hand, if the heating temperature exceeds 1200°C, austenite grain growth will be significant and the toughness of the base material will deteriorate. Therefore, the heating temperature should be 1200°C or lower. The heating temperature is also preferably 1190°C or lower, more preferably 1180°C or lower, even more preferably 1170°C or lower, and most preferably 1160°C or lower.

Hot Rolling Step In the method for producing a high-strength steel plate of the present invention, the hot rolling step is a step of hot rolling the steel material heated in the heating step, including one or more passes of rolling at 950°C or higher with a rolling reduction per pass of 10% or more and one or more passes of rolling at 900°C or lower with a rolling reduction per pass of 15% or more, under conditions where the cumulative rolling reduction at 900°C or lower is 50% or more, and the rolling finish temperature is between the _Ar3 temperature and 850°C.

The upper limit of the austenite non-recrystallization temperature range rises to around 900°C due to the inclusion of Nb. By rolling in the austenite non-recrystallization temperature range of 900°C or less, the austenite grains elongate, becoming finer in the thickness and width directions, and the density of dislocations within the grains introduced by rolling increases. Therefore, the cumulative reduction at 900°C or less is set to 50% or more. The cumulative reduction at 900°C or less is preferably 55% or more, more preferably 60% or more, even more preferably 65% or more, and most preferably 70% or more. There are no particular limitations on the upper limit of the cumulative reduction, but it is preferably 95% or less, more preferably 90% or less, and even more preferably 85% or less.

Furthermore, this effect is significantly exhibited by setting the cumulative reduction rate at 900°C or less to 50% or more and the rolling end temperature to be between the _Ar3 temperature and 850°C or less. Furthermore, in high-strength steel plates manufactured by hot rolling and cooling (described later), and in steel pipes made from such high-strength steel plates, strength, particularly strength in the medium temperature range, is improved. The _Ar3 temperature can be calculated using the following formula:
Ar ₃ (°C) = 910-310[C]-80[Mn]-20[Cu]-55[Ni]-15[Cr]-80[Mo]
In the formula, [C], [Mn], [Cu], [Ni], [Cr], and [Mo] respectively represent the contents (mass%) of C, Mn, Cu, Ni, Cr, and Mo in the base steel plate (high-strength steel plate). When some elements are not contained in the base steel plate, the content of the element may be set to "0" to determine the _Ar3 temperature.
Although there is no particular upper limit to the cumulative reduction rate, an excessively large cumulative reduction rate may impose an excessive load on the rolling mill, so the cumulative reduction rate is preferably 90% or less.

Furthermore, by including at least one pass of rolling at 950°C or higher with a reduction rate of 10% or more per pass, austenite grains are refined through recrystallization, improving the low-temperature toughness of the steel. Therefore, it is necessary to include at least one pass of rolling at 950°C or higher with a reduction rate of 10% or more per pass. It is preferable to include at least two passes of rolling at 950°C or higher with a reduction rate of 10% or more per pass. It is more preferable to include three or more passes of rolling at 950°C or higher with a reduction rate of 10% or more per pass, and even more preferable to include four or more passes. There is no particular upper limit to the number of passes for rolling at a reduction rate of 10% or more per pass, but it may be 50 passes or less, more preferably 30 passes or less, and even more preferably 15 passes or less. There is no particular upper limit to the reduction rate per pass at 950°C or higher, but it is preferably 40% or less.

Furthermore, by including at least one pass of rolling at 900°C or below with a reduction rate of 15% or more per pass, strain is introduced into the crystal grains, which promotes the refinement of austenite grains and improves the low-temperature toughness of the steel. Therefore, it is necessary to include at least one pass of rolling at 900°C or below with a reduction rate of 15% or more per pass. It is preferable to include at least two passes of rolling at 900°C or below with a reduction rate of 15% or more per pass. It is more preferable to include three or more passes of rolling at 900°C or below with a reduction rate of 15% or more per pass, and even more preferable to include four or more passes. There is no particular upper limit to the number of passes for rolling at a reduction rate of 15% or more per pass, but it may be 50 passes or less, more preferably 30 passes or less, and even more preferably 15 passes or less. There is no particular upper limit to the reduction rate per pass at 900°C or below, but it is preferable that it be 40% or less.

If the cumulative reduction rate at 900°C or less is less than 50% or the rolling end temperature exceeds 850°C, the austenite grains are not sufficiently refined, resulting in a deterioration in low-temperature toughness. Furthermore, the intragranular dislocation density decreases, resulting in a deterioration in strength in the intermediate temperature range. Furthermore, if the rolling end temperature is less than the _Ar3 temperature, a structure containing ferrite and bainite is formed, resulting in a deterioration in strength in the intermediate temperature range. Therefore, the cumulative reduction rate at 900°C or less is set to 50% or more, and the rolling end temperature is set to be equal to or higher than the _Ar3 temperature and equal to or higher than 850°C. The rolling end temperature is set to be equal to or higher than the _Ar3 temperature, preferably equal to or higher than the _Ar3 temperature + 5°C, more preferably equal to or higher than the _Ar3 temperature + 10°C, and even more preferably equal to or higher than the _Ar3 temperature + 15°C. Furthermore, the rolling end temperature is set to be equal to or lower than 850°C, preferably equal to or lower than 840°C, more preferably equal to or lower than 830°C, and even more preferably equal to or lower than 820°C.

Accelerated Cooling Step In the manufacturing method of a high-strength steel sheet of the present invention, the accelerated cooling step is a step of accelerated cooling of the hot-rolled steel sheet (hot-rolled sheet) obtained in the hot rolling step under the conditions of a cooling start temperature of _Ar3 temperature or higher, an average cooling rate of 5°C/s or higher, and a cooling stop temperature of 300 to 550°C. In this case, the average cooling rate refers to the average cooling rate from the cooling start temperature to the cooling stop temperature, and means the cooling rate obtained by dividing the difference between the cooling start temperature and the cooling stop temperature by the time required from the start of cooling to the stop of cooling.

In order to suppress the formation of ferrite on the front and back surfaces of the steel sheet and increase the bainite fraction, the cooling start temperature is set to be equal to or higher than the _Ar3 temperature. The cooling start temperature is preferably set to be equal to or higher than the _Ar3 temperature + ₁₀ °C, more preferably equal to or higher than the _Ar3 temperature + 15°C, even more preferably equal to or higher than the Ar3 temperature + 20°C, and most preferably equal to or higher than the _Ar3 temperature + 25°C. The upper limit is preferably set to be equal to or lower than the _Ar3 temperature + 150°C, more preferably equal to or lower than the _Ar3 temperature + 140°C, and even more preferably equal to or lower than the Ar3 temperature + 130°C, in order to prevent deterioration of low- _temperature toughness due to an increase in the average crystal grain size.

The strength of high-strength steel plate tends to increase as the average cooling rate during accelerated cooling increases. If the average cooling rate during accelerated cooling is less than 5°C/s, transformation begins at high temperatures, resulting in the formation of ferrite and pearlite in addition to bainite, and dislocation recovery also progresses during cooling. Therefore, if the average cooling rate is less than 5°C/s, sufficient strength cannot be achieved at room temperature or in the mid-temperature range. Furthermore, if the average cooling rate is less than 5°C/s, the effect of refining the structure is reduced, the crystal grain size does not decrease, and low-temperature toughness deteriorates. Therefore, the average cooling rate during accelerated cooling is set to 5°C/s or more. Furthermore, the average cooling rate during accelerated cooling is preferably 8°C/s or more, more preferably 10°C/s or more, even more preferably 15°C/s or more, and most preferably 20°C/s or more. While there is no particular upper limit to the cooling rate, to avoid an excessive increase in the martensite fraction, the average cooling rate is preferably 80°C/s or less, and more preferably 50°C/s or less.

Steel plate strength tends to increase as the accelerated cooling stop temperature decreases. If the accelerated cooling stop temperature exceeds 550°C, carbide growth is promoted and the amount of solute carbon decreases, resulting in insufficient strength, particularly in the mid-temperature range, in the cooled high-strength steel and in the steel pipe made from that high-strength steel. For this reason, the accelerated cooling stop temperature is set to 550°C or lower. It is preferably set to 540°C or lower, more preferably set to 530°C or lower, and even more preferably set to 520°C or lower. On the other hand, if the cooling stop temperature is lower than 300°C, the formation of low-temperature transformation products such as martensite, which has many mobile dislocations, becomes significant. As a result, long-term aging in the mid-temperature range promotes dislocation recovery, significantly reducing strength. Therefore, the accelerated cooling stop temperature is set to 300°C or higher. It is preferably set to 310°C or higher, more preferably set to 320°C or higher, even more preferably set to 330°C or higher, and most preferably set to 340°C or higher.

<Steel pipe manufacturing method>
The method for producing a steel pipe of the present invention includes a cold forming step and a welding step.

Cold forming process The cold forming process is a process of cold forming the high-strength steel plate of the present invention into a tubular shape. When manufacturing a steel pipe for transporting steam, the thickness of the steel plate is preferably 12 mm or more. Also, it is preferably 30 mm or less. More preferred ranges are the same as the plate thickness of the steel pipe described above. The method for cold forming the steel plate into a tubular shape is not particularly limited. Examples of forming methods include UOE forming, press bending, and roll forming.

Welding process The welding process is a process of welding the butt joints where the ends (one end and the other end) of the steel plates formed into a tubular shape in the cold forming process are butted together. The welding method is not particularly limited, but submerged arc welding or the like may be used for welding. It is preferable to expand the steel pipe after welding, as this improves the roundness of the pipe cross section. Heat treatment after steel pipe production may be performed depending on the desired properties, and is not particularly specified.

Steel plates (thickness: 15 to 25 mm) prepared under the manufacturing conditions shown in Tables 2-1 and 2-2 using steels A to AH having the chemical compositions shown in Table 1 were cold-formed, seam-welded, and then expanded by 1.0% to produce steel pipes with an outer diameter of 610 mm and a pipe thickness of 15 to 25 mm. Note that, in the manufacturing conditions shown in Tables 2-1 and 2-2, "rolling reduction" refers to the cumulative reduction at 900°C or less, "FT" refers to the rolling finish temperature, and "heat treatment" refers to long-term aging treatment. A 1.0% expansion ratio means that the inner diameter of the steel pipe is expanded by 1.0% using a pipe expander in the radial direction from the inner surface to the outer surface of the steel pipe.
Microstructure observation of the steel plate, evaluation of low-temperature toughness, tensile tests before and after aging, and tensile tests of the steel pipe before and after aging were carried out by the methods described below.

Observation of Steel Sheet Microstructure A sample for observing the steel structure was collected from the center of the sheet width of the steel sheet produced as described above, and the L cross section of the steel sheet (a cross section parallel to the rolling direction and parallel to the normal direction of the rolling surface) was mirror polished. After that, the crystal orientation of a randomly selected 1 mm x 1 mm region at the center of the sheet thickness was measured by electron backscatter diffraction (EBSD). Regions where the angle difference between adjacent pixels was 15° or more were considered to be grain boundaries, and image analysis was performed to determine the grain size of bainite.
The average crystal grain size darea (μm) was calculated from the area ai (μm ² ) occupied by each crystal grain and the circle-equivalent diameter di (μm) of each crystal grain using the following formula:
darea (μm) = Σ(ai・di)/Σai
The top 20% grain size of bainite represents the grain size of the grains that are 20% from the largest when the circle equivalent diameters of the grains are arranged in descending order of grain size and 20% of the total number of grains are selected from the largest.
Furthermore, the L-section of the steel sheet (a cross section parallel to the rolling direction and normal to the rolling surface) was mirror-polished and then subjected to nital etching to reveal the microstructure. Subsequently, using an optical microscope, the steel structure was photographed in five randomly selected fields of view (7.1 × ^{10-2 mm2} , magnification: 400x) at the center of the sheet thickness. The bainite fraction, ferrite fraction, and island martensite fraction in the photographed fields were measured as area fractions using an image analyzer (Fiji), and the bainite grain size was measured using an image analyzer (OIM Analysis, manufactured by TSL). The average values of the five fields were used as the bainite fraction, ferrite fraction, island martensite fraction, and bainite grain size.

Furthermore, as described above, the number of precipitates was determined using the evaluation method used in the embodiment.

Evaluation of Low-Temperature Toughness of Steel Plates Furthermore, DWTT in accordance with API 5L was performed to evaluate the low-temperature toughness of the steel plates. DWTT test specimens were taken so that the thickness direction of the steel plate coincided with the thickness direction of the test specimen and the width direction of the steel plate coincided with the longitudinal direction of the test specimen. A test specimen with a ductile fracture area ratio of 85% or more and a fracture transition temperature of -40°C or less was evaluated as good.

Tensile test of steel sheet before and after aging For the steel sheet properties (before aging), tensile test specimens were taken so that the longitudinal direction of the tensile test specimen was perpendicular to the rolling direction of the steel sheet and perpendicular to the sheet thickness direction, and a tensile test was carried out at 350°C to determine the yield strength. In the tensile test, a round bar test specimen with a diameter of 6 mm was used, and the crosshead speed was 0.15 mm/min. A yield strength of 555 MPa or more in the tensile test at 350°C was evaluated as good.

To evaluate the properties after long-term aging (heat treatment at 400°C for 2335 hours), tensile test specimens were taken so that the longitudinal direction of the test specimen was perpendicular to the steel plate rolling direction and perpendicular to the plate thickness direction, and heat treatment was carried out in an N (nitrogen) atmosphere furnace at 400°C for 2335 hours. A tensile test was then carried out at 350°C to evaluate the yield strength. The tensile test was carried out using a round bar test specimen with a diameter of 6 mm, with a crosshead speed of 0.15 mm/min. A yield strength of 555 MPa or more in the tensile test at 350°C was evaluated as good.

In addition, the change in yield strength when held for a long time in the medium temperature range was evaluated by measuring the difference in yield strength before and after the long-term aging described above, i.e., (yield strength before aging - yield strength after aging).

Furthermore, to assess the properties after long-term aging (heat treatment at 400°C for 2,335 hours), 6 mm diameter round bar test specimens were taken from steel pipes with the circumferential direction of the pipe aligned with the longitudinal direction of the test specimen. These were then heat treated in a furnace at 400°C for 2,335 hours in an N (nitrogen) atmosphere, and tensile tests were conducted at 350°C. The yield strength was evaluated in the same way as for steel plate, with a yield strength of 555 MPa or more at 350°C being rated as good.

As mentioned above, Tables 2-1 and 2-2 show the manufacturing conditions of the steel plates and the test results of the steel plates and steel pipes.
Inventive Examples (Nos. 1 to 29) in which the chemical composition and steel plate manufacturing conditions were both within the ranges of the present invention, the steel plates and steel pipes had yield strengths of 555 MPa or more at 350°C, and also obtained yield strengths of 555 MPa or more at 350°C after long-term aging of the steel plates and steel pipes, with (yield strength before aging - yield strength after aging) being 50 MPa or less, and the ductile fracture area ratio: DWTTSA-40°C obtained by DWTT (test temperature: -40°C) in accordance with API 5L being 85% or more, and a fracture appearance transition temperature of -40°C or less was obtained.

On the other hand, the comparative examples (Nos. 30 to 59), whose chemical compositions or steel sheet manufacturing conditions were outside the scope of the present invention, were inferior to the examples of the present invention in at least one of the following: yield strength at 350°C and/or yield strength at 350°C before and after long-term aging of 555 MPa or more, ductile fracture area ratio obtained by DWTT (test temperature: -40°C) of 85% or more, and fracture appearance transition temperature of -40°C or less.

 

 

Claims (6)

The component composition is, in mass%,
C: 0.04-0.09%,
Si: 0.03-0.25%,
Mn: 1.5-2.5%,
P: 0.020% or less,
S: 0.002% or less,
Mo: 0.10-0.50%,
Nb: 0.010-0.055%,
Ti: 0.005 to 0.020%,
Ca: 0.0040% or less,
Al: 0.01-0.04%,
N: 0.006% or less, the balance being Fe and unavoidable impurities;
X (%) represented by formula (1) is 0.65% or more,
the microstructure has bainite in an area ratio of 80% or more at the center of the plate thickness, the average grain size of the bainite is 25 μm or less, and the minimum grain size of the top 20% of the bainite grains with the largest grain size is 60 μm or less, the toughness of the steel plate is such that the ductile fracture area ratio obtained by DWTT at -40°C is 85% or more,
A high-strength steel plate having a yield strength of 555 MPa or more before and after aging under the condition of Larson Miller Parameter LMP = 15700, as defined by the following formula (2).
X (%) = 0.35Cr + 0.9Mo + 12.5Nb + 8V... (1)
The element symbols in formula (1) represent the content (mass%) of each element. Elements that are not contained are substituted with 0.
LMP=(T+273)×(20+log(t))...(2)
T: Heat treatment temperature (°C)
t: heat treatment time (hours) The component composition is further expressed in mass % as follows:
Cr: 0.50% or less,
V: 0.070% or less,
Cu: 0.50% or less,
2. The high-strength steel plate according to claim 1, further comprising one or more of the following: Ni: 0.50% or less. Furthermore, the sheet has an area ratio of ferrite of 10% or less and an area ratio of island martensite of 10% or less at the center of the sheet thickness,
the total number of Ti-based precipitates having a diameter of 100 nm or less, Nb-based precipitates having a diameter of 100 nm or less, V-based precipitates having a diameter of 100 nm or less, Mo-based precipitates having a diameter of 100 nm or less, Cr-based precipitates having a diameter of 100 nm or less, Al-based precipitates having a diameter of 100 nm or less, and composite precipitates having a diameter of 100 nm or less containing two or more elements selected from Ti, Nb, V, Mo, Cr, and Al is 50,000 to 1,000,000 ^per 1 mm2 of the test area, at the center of the sheet thickness;
3. The high-strength steel plate according to claim 1, wherein a difference in yield strength before and after aging, obtained by subtracting the yield strength after aging from the yield strength before aging, is 50 MPa or less. A steel pipe made using the high-strength steel plate described in any one of claims 1 to 3. The method for producing a high-strength steel plate according to any one of claims 1 to 3, comprising: a heating step of heating a steel material to 1000 to 1200 ° C;
a hot rolling process in which the steel material heated in the heating process is hot rolled under the conditions of one or more passes of rolling at 950°C or higher with a rolling reduction of 10% or more per pass and one or more passes of rolling at 900°C or lower with a rolling reduction of 15% or more per pass, a cumulative rolling reduction of 50% or more at 900°C or lower, and a rolling end temperature of _Ar3 temperature or higher and 850°C or lower;
and an accelerated cooling step of accelerated cooling the hot-rolled steel sheet obtained in the hot rolling step under conditions of a cooling start temperature of Ar ₃ temperature or higher, an average cooling rate of 5°C/s or higher, and a cooling stop temperature of 300 to 550°C. a cold forming step of cold forming the high strength steel plate according to any one of claims 1 to 3 into a tubular shape;
a welding process for welding the butt joints at the butt joints of the ends of the steel plates formed into a tubular shape in the cold forming process;
A method for manufacturing a steel pipe having the above structure.