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

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

[0001] 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 intermediate temperature range, a method for manufacturing the same, and a steel pipe made of the high-strength steel plate and a method for manufacturing the same. In particular, the high-strength steel plate of the present invention can be suitably used as a material for high-strength steel pipes for steam piping.

[0002] 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, 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, increasing energy demand has led to demand for larger diameter and higher strength steel pipes to improve heavy oil recovery rates and reduce construction costs. Furthermore, when laying these steel pipes, construction work and hydraulic pressure testing are typically carried out in the early spring when temperatures are high. However, due to increasing demand, superior low-temperature toughness is also required to enable construction work in the cold winter months.

[0003] Patent Documents 1 and 2 are prior art examples of 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 the outer diameter of this seamless steel pipe being a maximum of 16 inches. It is difficult to further increase the diameter of seamless steel pipes. Furthermore, in order to obtain a strength equal to or greater than API X80 grade, the seamless steel pipe requires the addition of large amounts of alloying elements.

[0004] Incidentally, Patent Documents 3 and 4 disclose manufacturing techniques 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 techniques for high-strength steel pipes that are manufactured by TMCP (thermo-mechanical control process) and have strength of API X80 grade or higher.

[0005] Japanese Patent Publication No. 2000-290728 Japanese Patent No. 4821939 Japanese Patent No. 5055736 Japanese Patent No. 4741528 Japanese Patent No. 6241569

[0006] The steel pipe manufactured by the steel pipe manufacturing method disclosed in Patent Document 3 has high temperature properties in the medium temperature range that satisfy API X80 grade. However, Patent Document 3 does not disclose the strength properties of the steel pipe when used for a long period of time in the medium temperature range.

[0007] Patent Document 4 describes the creep characteristics of steel pipes after long-term aging, and although the developed steel shows an improvement in fracture strength, the safety criteria are not clearly defined. Furthermore, fracture strength values ​​are occasionally found to be 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.

[0008] The steel pipe disclosed in Patent Document 5 satisfies the API X80 grade in terms of high temperature properties in the intermediate temperature range, but has a small LMP, and no consideration is given to the strength properties when the steel pipe is used for a long period of time, nor to low temperature toughness.

[0009] As described above, the prior art has been unable to obtain a high-strength steel pipe for steam piping that satisfies all of the following requirements: a large diameter, strength characteristics required for high-strength steel pipe for steam transportation, and excellent low-temperature toughness.

[0010] The present invention has been made to solve the above-mentioned problems, and has an object to provide a high-strength steel plate that can be used as a raw material for steel pipes that satisfy all of the requirements for large diameter, excellent low-temperature toughness, and strength characteristics required for high-strength steel pipes for steam transportation, and a method for manufacturing the same. Another object of the present invention is to provide a steel pipe made from the high-strength steel plate and a method for manufacturing the same.

[0011] 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 in the DWTT is set to -40°C in anticipation of a decrease in toughness due to work hardening during pipe making.

[0012] The present inventors have conducted extensive research into the properties of high-strength steel plates for large-diameter steel pipes in the medium temperature range, and have found that by appropriately selecting the chemical composition and manufacturing conditions, it is possible to obtain high-strength steel plates that can be used to manufacture high-strength steel pipes that have the strength characteristics and low-temperature toughness required for high-strength steel pipes for transporting steam, even though they are large in diameter.

[0013] The present invention was made based on the above findings and further studies, and is comprised of the following: [1] A component composition, in mass %, containing C: 0.04 to 0.09%, Si: 0.03 to 0.25%, Mn: 1.5 to 2.5%, P: 0.020% or less, S: 0.002% or less, Mo: 0.10 to 0.50%, Nb: 0.010 to 0.055%, Ti: 0.005 to 0.020%, Ca: 0.0040% or less, Al: 0.01 to 0.04%, N: 0.006% or less, with the balance being Fe and unavoidable impurities, and X (%) represented by formula (1) is 0.65% or more, A high-strength steel sheet having a microstructure having bainite in an area fraction of 80% or more at the center of the sheet thickness, the bainite having an average grain size of 25 μm or less, and the smallest grain size of the top 20% of the largest grains of the bainite being 60 μm or less, the toughness of the steel sheet being such that the ductile fracture area ratio obtained by DWTT at -40°C is 85% or more, and the yield strength before and after aging under the Larson Miller Parameter (LMP) of 15700, as defined by the following formula (2), is 555 MPa or more: X (%) = 0.35Cr + 0.9Mo + 12.5Nb + 8V (1) The element symbols in formula (1) represent the content (mass%) of each element. Also, 0 is substituted for elements that are not contained. LMP = (T + 273) × (20 + log(t)) (2), T: heat treatment temperature (°C), t: heat treatment time (hours) [2] The high-strength steel plate according to [1], wherein the chemical composition further contains, in mass%, one or more of Cr: 0.50% or less, V: 0.070% or less, Cu: 0.50% or less, and Ni: 0.50% or less. [3] Furthermore, the steel 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, and the total 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 containing two or more elements selected from Ti, Nb, V, Mo, Cr, and Al is 1 mm or less in an inspection area. ２1. The high-strength steel plate according to claim 1 or 2, wherein 50,000 to 1,000,000 particles are present per 1000 kJ / cm2, and 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 of [1] to [3]. [5] A method for producing a high-strength steel plate according to any of [1] to [3], comprising: a heating step of heating a steel material to 1000 to 1200°C; and one or more passes of rolling the steel material heated in the heating step 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, wherein the cumulative rolling reduction at 900°C or lower is 50% or more and the rolling end temperature is Ar. ３ a hot rolling step of hot rolling the hot rolled steel sheet under the condition of a temperature of 850°C or higher and a cooling start temperature of Ar ３ and an accelerated cooling step of performing accelerated cooling under conditions of a temperature above 300°C, an average cooling rate of 5°C / s or more, and a cooling stop temperature of 300 to 550°C. [6] A method for producing a steel pipe, comprising: a cold forming step of cold-forming the high-strength steel plate according to any one of [1] to [3] into a tubular shape; and a welding step of butting together the ends of the steel plate formed into a tubular shape in the cold forming step, and welding the butt joints.

[0014] According to the present invention, it is possible to obtain a high-strength steel pipe that has the strength and toughness characteristics required for a high-strength steel pipe for transporting steam, even though it has a large diameter.

[0015] Hereinafter, embodiments of the present invention will be described, but the present invention is not limited to the following embodiments.

[0016] <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."

[0017] C: 0.04 to 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 a deterioration in toughness and weldability. For this reason, the upper limit of the C content is set to 0.09%. Furthermore, the C content is preferably 0.08% or less, and more preferably 0.07% or less.

[0018] Si: 0.03 to 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. For this reason, 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.

[0019] Mn: 1.5 to 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.

[0020] 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.

[0021] 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.

[0022] Mo: 0.10 to 0.50% Mo solid-solution strengthens steel and improves its hardenability. Solid-solution strengthening and improved hardenability increase strength, particularly in the mid-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.

[0023] Nb: 0.010 to 0.055% Nb suppresses grain growth during slab heating and rolling. This refines the microstructure, imparting sufficient strength and toughness to the steel. Nb is also an essential component for forming carbides to ensure strength in the mid-temperature range. This effect is most pronounced at an Nb content of 0.010% or more. For this reason, 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 0.053% or less, more preferably 0.050% or less, further preferably 0.048% or less, and most preferably 0.045% or less.

[0024] Ti: 0.005 to 0.020% Ti forms TiN, which inhibits grain growth during slab heating and in weld heat-affected zones, refines the microstructure, and improves toughness. 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.

[0025] 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.

[0026] Al: 0.01 to 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 preferably 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.

[0027] 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.

[0028] 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 such cases, 0 can be substituted for "Cr" or "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. In addition, 0 is substituted for elements that are not included.

[0029] The high strength steel sheet of the present invention may further contain one or more of Cr, V, Cu and Ni for the purpose of further improving the properties.

[0030] 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.

[0031] 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.

[0032] 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.

[0033] 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.

[0034] In the present invention, the Cu content, Ni content, Cr content, and Mo content are preferably adjusted so that Cu + Ni + Cr + Mo (element symbols represent the content (mass%) of each element) is 1.00% or less. In order to achieve both the effect of increasing strength by these elements and production 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 restriction on the 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.

[0035] The remainder other than the above components is Fe and inevitable impurities. These impurities are inevitably mixed in from raw materials, the manufacturing process, or manufacturing equipment, and are allowed to be present to the extent that they do not impair the object of the present invention. Examples of raw materials include iron ore, reduced iron, and scrap. Examples of inevitable impurities include Pb, Zn, Sn, As, B, Sb, Bi, Co, H, O, and REM.

[0036] Next, the structure of the high strength steel sheet of the present invention will be described.

[0037] 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.

[0038] Ferrite: Area Fraction of 10% or Less (Preferred 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 lower limit for the area fraction of ferrite, but from the viewpoint of improving deformability, the area fraction of ferrite is preferably 1% or more.

[0039] Island martensite (MA): area fraction 10% or less (preferred condition) Island martensite (MA: martensite-austenite constituent) is a very hard phase and may act as a fracture initiation site, 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 an area fraction of 1% or more is preferred, and 2% or more is more preferred. Furthermore, 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.

[0040] 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.

[0041] Furthermore, by refining the average grain size, low-temperature toughness is improved. ３When cooling is initiated around the temperature, there is a limit to how much the average grain size can be refined. In the present invention, it is essential to suppress the formation of even larger grains. Coarse bainite is highly likely to become the origin of fracture, and if the top 20% of the largest grains among all bainite grains have a large grain size, low-temperature toughness deteriorates. In particular, if the minimum grain size of the top 20% of the largest grains exceeds 60 μm, they are likely to become the origin of fracture. Therefore, the minimum grain size of the top 20% of the largest grains 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. There is no particular limitation on the lower limit, but 10 μm or more is preferred, and 15 μm or more is more preferred.

[0042] Here, the average grain size of bainite is determined as follows: After mirror polishing the L-section of the steel sheet (a section parallel to the rolling direction and parallel to the normal direction of the rolling surface), a randomly selected 1 mm × 1 mm ２ The crystal orientation of the region was measured by electron backscatter diffraction (EBSD), and the region where the angle difference between adjacent pixels was 15° or more was determined as a grain boundary by image analysis. The EBSD measurement conditions were an acceleration voltage of 17 kV and a measurement pitch of 0.8 μm. The average grain size darea (μm) was calculated by dividing the area ai (μm) occupied by each grain by ２ ) 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 with the largest grain size represents the grain size of the smallest grain when the circle-equivalent diameter of each grain is arranged in descending order of grain size and 20% of the total number of grains are selected from the largest.

[0043] The total 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 within an inspection area of ​​1 mm. ２ The reasons for limiting the size 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, the number of precipitates formed is limited to 1 mm. ２ If the number of precipitates is less than 50,000 per mm, the effect of suppressing austenite grain growth is weak, and the decrease in strength in the medium temperature range may not be suppressed. ２ It is preferable that the number of precipitates having a diameter of 100 nm or less is 50,000 or more per mm. 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. ２ If the number of particles exceeds 1,000,000 per mm, the fine precipitates will aggregate and coarsen, which may weaken the effect of inhibiting austenite grain growth and result in a significant decrease in strength in the intermediate temperature range. ２ It is preferable that the number of molecules per molecule is 1,000,000 or less, more preferably 950,000 or less, even more preferably 900,000 or less, and most preferably 850,000 or less.

[0044] Here, the method for measuring the density and size (diameter) of the precipitates will be explained. 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 of 100 nm or less is counted by counting at least 100 μm. ２ For example, when observing at a magnification of 100,000 times with one visual field measuring 2000 nm x 2000 nm, the observed area per visual field is 4 μm ２ Therefore, 25 visual fields are randomly observed. The measurement results are converted into the number per unit area. ２ ) and there are five, the density of the precipitates is 1 mm ２ The density of the precipitates can be calculated as 50,000 particles per 1 mm2 of the test area. ２ The diameter of the precipitate is the average value of the major axis (long side) and minor axis (short side).

[0045] 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 Larson Miller Parameter LMP = 15700. Aging treatment at LMP = 15700 refers to aging treatment performed under conditions of heat treatment temperature and heat treatment time such that the LMP, expressed by the following formula (2), becomes 15700. LMP = 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) = 15700 refers to a Larson Miller Parameter (LMP) of 15650 or more 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 plate 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 refers to 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. There is no particular upper limit, but the yield strength is preferably 840 MPa or less. Note that, as described in the Examples, in the present invention, the yield strengths of test specimens taken from both the steel plate and the steel pipe are 555 MPa or more before and after the aging treatment.

[0046] 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 (preferred 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 aging is 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 period of 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. For this reason, 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. Furthermore, an example of the aging treatment conditions under the above conditions is heat treatment at 400°C for 2335 hours.

[0047] 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.

[0048] <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 a high-strength welded steel pipe for steam transportation, even when it has a large diameter.

[0049] "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, the diameter can be increased while maintaining the strength characteristics required for high-strength welded steel pipes for steam transportation.

[0050] The thickness of the steel pipe is not particularly limited, but is 12 to 30 mm in the case of steam transportation. 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.

[0051] The strength characteristics required for high-strength welded steel pipe for steam transportation are, like the above-mentioned high-strength steel, a yield strength of 555 MPa or more at 350°C before and after aging under the condition of Larson Miller Parameter (LMP) = 15700. It 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. There is no particular upper limit, but the yield strength at 350°C before and after the aging is preferably 840 MPa or less, more preferably 830 MPa or less.

[0052] <Method for 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, the temperature refers 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.

[0053] 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.

[0054] In the hot rolling process described below, in order to sufficiently promote austenitization and solid solution of carbides and obtain sufficient strength at room temperature and in the intermediate temperature range, the heating temperature of the steel material needs to 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 becomes significant and the toughness of the base material deteriorates. Therefore, the heating temperature is set to 1200°C or lower. Furthermore, the heating temperature is 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.

[0055] In the method for producing a high-strength steel plate of the present invention, the hot rolling step includes rolling the steel material heated in the heating step at 950°C or higher at a rolling reduction rate of 10% or more per pass at 950°C or higher and at 900°C or lower at a rolling reduction rate of 15% or more per pass at a rolling end temperature of Ar. ３ This is a process of hot rolling under the condition of a temperature of 850°C or higher and 850°C or lower.

[0056] 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 dislocation density 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. The upper limit of the cumulative reduction is not particularly limited, but is preferably 95% or less, more preferably 90% or less, and even more preferably 85% or less.

[0057] Furthermore, the cumulative reduction rate at 900°C or less is 50% or more, and the rolling end temperature is Ar ３ This effect is particularly pronounced when the temperature is set to 850°C or higher. Furthermore, in a high-strength steel plate manufactured by hot rolling and cooling (described later), and in a steel pipe made of the high-strength steel plate, the strength, particularly in the intermediate temperature range, is improved. ３ The temperature can be calculated by the following formula: Ar ３ (°C) = 910 - 310 [C] - 80 [Mn] - 20 [Cu] - 55 [Ni] - 15 [Cr] - 80 [Mo] where [C], [Mn], [Cu], [Ni], [Cr] and [Mo] respectively mean the contents (mass%) of C, Mn, Cu, Ni, Cr and Mo in the steel of the base steel plate (high strength steel plate). In addition, when some elements are not contained in the base steel plate, the content of the element is set to "0" and Ar is used. ３ 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, and therefore the cumulative reduction rate is preferably 90% or less.

[0058] Furthermore, by including one or more passes of rolling at 950°C or higher with a reduction rate of 10% or more per pass, austenite grains are refined by recrystallization, thereby improving the low-temperature toughness of the steel. Therefore, it is necessary to include one or more passes of rolling at 950°C or higher with a reduction rate of 10% or more per pass. It is also preferable to include two or more 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. The upper limit of the number of passes of rolling at a reduction rate of 10% or more per pass is not particularly limited, but may be 50 passes or less, more preferably 30 passes or less, and even more preferably 15 passes or less. The upper limit of the reduction rate per pass at 950°C or higher is not particularly limited, but is preferably 40% or less.

[0059] Furthermore, by including one or more passes of rolling at 900°C or less with a reduction rate of 15% or more per pass, strain is introduced into the crystal grains, thereby promoting the refinement of austenite grains and improving the low-temperature toughness of the steel. Therefore, it is necessary to include one or more passes of rolling at 900°C or less with a reduction rate of 15% or more per pass. It is preferable to include two or more passes of rolling at 900°C or less 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 less with a reduction rate of 15% or more per pass, and even more preferable to include four or more passes. The upper limit of the number of passes of rolling at a reduction rate of 15% or more per pass is not particularly limited, but may be 50 passes or less, more preferably 30 passes or less, and even more preferably 15 passes or less. The upper limit of the reduction rate per pass at 900°C or less is not particularly limited, but is preferably 40% or less.

[0060] 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 refined sufficiently, and low-temperature toughness deteriorates. In addition, the dislocation density within the grains decreases, and strength in the intermediate temperature range deteriorates. Furthermore, if the rolling end temperature is Ar ３If the rolling temperature is less than 900°C, a structure containing ferrite and bainite is formed, and the strength in the intermediate temperature range deteriorates. Therefore, the cumulative reduction rate at 900°C or less is set to 50% or more, and the rolling end temperature is set to Ar. ３ The rolling end temperature is Ar ３ temperature or higher, preferably Ar ３ The temperature is set to +5°C or higher, and more preferably, Ar ３ The temperature is set to +10°C or higher, and more preferably, Ar ３ The rolling end temperature is set to 850°C or lower, preferably 840°C or lower, more preferably 830°C or lower, and even more preferably 820°C or lower.

[0061] In the method for producing a high-strength steel sheet according to the present invention, the accelerated cooling step is a step of cooling the hot-rolled steel sheet (hot-rolled sheet) obtained in the hot rolling step at a cooling start temperature of Ar. ３ This is a process of accelerated cooling under the conditions of a 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.

[0062] In order to suppress the generation of ferrite on the front and back surfaces of the steel sheet and increase the bainite fraction, the cooling start temperature is set to Ar. ３ The cooling start temperature is Ar ３ The temperature is preferably +10°C or higher, and Ar ３ It is more preferable to set the temperature at +15°C or higher, and Ar ３ It is more preferable to set the temperature at +20°C or higher, and Ar ３ The most preferable temperature is +25°C or more. The upper limit is Ar to prevent deterioration of low-temperature toughness due to an increase in the average grain size. ３ The temperature is preferably +150°C or less, and Ar ３ It is more preferable to set the temperature at +140°C or less, and Ar ３ It is more preferable to set the temperature at +130° C. or lower.

[0063] The strength of high-strength steel plates tends to increase with increasing average cooling rate during accelerated cooling. When 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, when the average cooling rate is less than 5°C / s, sufficient strength cannot be obtained at room temperature or in the intermediate temperature range. Furthermore, when 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, more preferably 50°C / s or less.

[0064] The strength of steel sheets tends to increase as the cooling stop temperature of accelerated cooling decreases. If the cooling stop temperature of accelerated cooling exceeds 550°C, carbide growth is promoted and the amount of solute carbon decreases, resulting in insufficient strength, particularly in the intermediate temperature range, in the cooled high-strength steel and in the steel pipe made of that high-strength steel. Therefore, the cooling stop temperature of accelerated cooling 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 intermediate temperature range promotes dislocation recovery, significantly reducing strength. Therefore, the cooling stop temperature of accelerated cooling 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.

[0065] <Method for Producing Steel Pipe> The method for producing a steel pipe of the present invention includes a cold forming step and a welding step.

[0066] 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.

[0067] 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.

[0068] Steel plates (15-25 mm thick) 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, followed by seam welding and subsequent 1.0% expansion to produce steel pipes with an outer diameter of 610 mm and a pipe thickness of 15-25 mm. 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 was expanded by 1.0% in the radial direction from the inner surface to the outer surface using a pipe expander. Microstructure observation of the steel plates, low-temperature toughness evaluation, tensile tests before and after aging, and tensile tests of the steel pipes before and after aging were performed using the methods described below.

[0069] Observation of the Microstructure of Steel Sheets A sample for observing the steel structure was taken from the center of the sheet width of the steel sheet produced as described above, and the L-section of the steel sheet (a 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 regarded as grain boundaries, and image analysis was performed to determine the grain size of bainite. The average grain size darea (μm) was calculated by the area ai (μm) occupied by each grain. ２ ) 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 20% largest grains when the circle equivalent diameter of each crystal grain is arranged in descending order of grain size, and 20% of the total number of crystal grains are selected from the largest. 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. Then, using an optical microscope, randomly selected 7.1 x 10 －２ mm ２ The steel structure was photographed in five fields of view (magnification: 400x), and the bainite fraction, ferrite fraction, and island martensite fraction in the photographed fields were measured as area percentages 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 of view were taken as the bainite fraction, ferrite fraction, island martensite fraction, and bainite grain size.

[0070] The number of precipitates was determined using the evaluation method used in the embodiment, as described above.

[0071] 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.

[0072] 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.

[0073] To evaluate the properties after long-term aging (heat treatment at 400°C for 2335 hours), tensile test specimens were taken so that the direction perpendicular to the rolling direction of the steel sheet and perpendicular to the sheet thickness direction was the longitudinal direction of the test specimen, and heat treatment was carried out at 400°C for 2335 hours in an N (nitrogen) atmosphere furnace, followed by a tensile test at 350°C to evaluate the yield strength. In the tensile test, a round bar test specimen with a diameter of 6 mm was used, and a crosshead speed of 0.15 mm / min was used. A yield strength of 555 MPa or more in the tensile test at 350°C was evaluated as good.

[0074] Furthermore, the change in yield strength when held in the intermediate temperature range for a long time was evaluated by measuring the difference in yield strength before and after the long-term aging, that is, (yield strength before aging - yield strength after aging).

[0075] Furthermore, to evaluate the properties after long-term aging (heat treatment at 400°C for 2335 hours), round bar test specimens with a diameter of 6 mm were also taken from steel pipes, with the circumferential direction of the pipe being the longitudinal direction of the test specimen, and heat treatment was carried out in a furnace at 400°C for 2335 hours in an N (nitrogen) atmosphere furnace. Tensile tests were then carried out at 350°C, and the yield strength was evaluated in the same manner as for steel plates, with yield strengths of 555 MPa or more at 350°C being evaluated as good.

[0076] As described 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. In the invention examples (Nos. 1 to 29) in which the chemical composition and manufacturing conditions of the steel plates were within the ranges of the invention, the steel plates and steel pipes had a yield strength of 555 MPa or more at 350°C, and also obtained a yield strength 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.

[0077] 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: a yield strength at 350°C and / or a yield strength at 350°C before and after long-term aging of 555 MPa or more, a ductile fracture area ratio obtained by DWTT (test temperature: -40°C) of 85% or more, and a fracture appearance transition temperature of -40°C or less.

[0078]

[0079]

[0080]

Claims

1. The chemical composition, in mass%, is as follows: C: 0.04 to 0.09%, Si: 0.03 to 0.25%, Mn: 1.5 to 2.5%, P: 0.020% or less, S: 0.002% or less, Mo: 0.10 to 0.50%, Nb: 0.010 to 0.055%, Ti: 0.005 to 0.020%, Ca: 0.0040% or less, Al: 0.01 to 0.04%, N: 0.006% or less, with the balance being Fe and unavoidable impurities; and X (%) represented by formula (1) is 0.65% or more. A high-strength steel sheet having a microstructure having bainite in an area fraction of 80% or more at the center of the sheet thickness, the bainite having an average grain size of 25 μm or less, and the smallest grain size of the top 20% of the largest grains of the bainite being 60 μm or less, the toughness of the steel sheet being such that the ductile fracture area ratio obtained by DWTT at -40°C is 85% or more, and the yield strength before and after aging under the Larson Miller Parameter (LMP) of 15700, as defined by the following formula (2), is 555 MPa or more: X (%) = 0.35Cr + 0.9Mo + 12.5Nb + 8V (1) The element symbols in formula (1) represent the content (mass%) of each element. Also, 0 is substituted for elements that are not contained. LMP = (T + 273) × (20 + log (t)) (2) T: heat treatment temperature (°C) t: heat treatment time (hours) 2. A high-strength steel plate according to claim 1, wherein the chemical composition further contains, in mass %, one or more of the following: Cr: 0.50% or less, V: 0.070% or less, Cu: 0.50% or less, and Ni: 0.50% or less.

3. 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, and the total of Ti-based precipitates of 100 nm or less in diameter, Nb-based precipitates of 100 nm or less in diameter, V-based precipitates of 100 nm or less in diameter, Mo-based precipitates of 100 nm or less in diameter, Cr-based precipitates of 100 nm or less in diameter, Al-based precipitates of 100 nm or less in diameter, and composite precipitates of 100 nm or less containing two or more elements selected from Ti, Nb, V, Mo, Cr, and Al is 1 mm or less in diameter. ２ 3. The high-strength steel plate according to claim 1, wherein the number of such defects is 50,000 to 1,000,000 per one grain, and 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 made from the high-strength steel plate according to any one of claims 1 to 3.

5. A method for producing 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; and rolling the steel material heated in the heating step at least one pass at 950°C or higher with a rolling reduction of 10% or more per pass and at 900°C or lower with a rolling reduction of 15% or more per pass, wherein the cumulative rolling reduction at 900°C or lower is 50% or more and the rolling end temperature is Ar. ３ a hot rolling step of hot rolling the hot rolled steel sheet under the condition of a temperature of 850°C or higher and a cooling start temperature of Ar ３ and an accelerated cooling step of performing accelerated cooling under conditions of a 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 method for manufacturing a steel pipe, comprising: a cold forming step of cold-forming a high-strength steel plate according to any one of claims 1 to 3 into a tubular shape; and a welding step of butting together the ends of the steel plate formed into a tubular shape in the cold forming step, and welding the butt joints.

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