Classifications

• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese

Definitions

• the present invention relates to a steel plate having excellent low-temperature toughness and a tensile strength of 620 MPa or more after long-term aging in the medium temperature range, a method for manufacturing the same, and a steel pipe made from the steel plate and a method for manufacturing the same.
• the 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 disclose methods for manufacturing 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, but the maximum outer diameter of these seamless steel pipes is 16 inches, making it difficult to further increase the diameter of seamless steel pipes.
• Patent Documents 3, 4, and 5 disclose manufacturing techniques for high-strength steel pipes that are manufactured by welding and can be made larger in diameter, and have a strength of API X80 or higher.
• Patent Document 3 Steel pipes manufactured using the steel pipe manufacturing method disclosed in Patent Document 3 have high-temperature properties in the medium temperature range that meet API X80 grade standards. 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 discloses a manufacturing technique for high-strength steel plate with a yield strength of 700 MPa or more.
• a manufacturing technique for high-strength steel plate with a yield strength of 700 MPa or more in order to ensure strength in the medium temperature range using the high-strength steel plate manufacturing technique disclosed in Patent Document 4, it is necessary to add large amounts of alloying elements.
• the tensile strength of the high-strength steel plate described in Patent Document 4 decreases significantly when held in the medium temperature range for long periods of time.
• Steel pipes of API X80 or higher are required to meet the requirements of a yield strength of 555 MPa or more and a tensile strength of 620 MPa or more in a tensile test. Furthermore, other required properties for high-strength steel pipes for steam transport include meeting the above strength characteristics even before and after long-term aging in the medium temperature range, as well as a large diameter, excellent low-temperature toughness, and excellent toughness in the heat-affected zone formed during welding.
• the present invention has been made to solve the above-mentioned problems, and the steel plate that serves as the material for the steel pipe must satisfy all of the following requirements: a yield strength of 555 MPa or more both before and after aging, a tensile strength of 620 MPa or more, as determined by a tensile test at 350°C; excellent low-temperature toughness; and excellent toughness in the weld heat-affected zone formed during welding.
• the present invention also aims to provide a steel plate that serves as the material for the steel pipe and a method for manufacturing the same.
• Another object of the present invention is to provide a steel pipe made from the above-mentioned steel plate and a method for manufacturing the same.
• 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.
• excellent toughness in the weld heat affected zone formed during welding means that the absorbed energy (vE -40 ) at -40°C is 60 J or more in a Charpy impact test using a test piece taken from the weld heat affected zone.
• large diameter means that the outer diameter of the steel pipe is 400 mm or more.
• the present inventors have conducted extensive research into the properties of 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 steel plates that can be used to manufacture steel pipes that have the strength properties and low-temperature toughness required for high-strength steel pipes for transporting steam, even though they are large diameter pipes.
• Nb carbides and V carbides refer to simple carbides containing Nb or V, and composite carbides containing Nb and V.
• P eff (%) (0.13Nb+0.24V-0.125Ti)/(C+0.86N)...(1)
• the element symbols in formula (1) represent the content (mass%) of each element. Elements that are not contained are substituted with 0.
• Nb and V are elements that form carbides in steel. Strengthening steel by the precipitation of NbC has been a common practice. V-based carbides are also resistant to aggregation and coarsening even when held at high temperatures for long periods of time, making them useful elements for ensuring high-temperature creep strength.
• the heating rate (temperature increase rate) during reheating after accelerated cooling is increased to suppress the growth of precipitates during heating. This allows a large amount of fine precipitates containing Nb or carbides containing Nb and V to precipitate in the steel, thereby suppressing strength degradation in the medium temperature range.
• the hot-rolled steel sheet in the reheating after accelerated cooling, is heated in an atmospheric furnace at a heating rate faster than that conventionally adopted in industry, thereby suppressing the growth of carbides containing Nb or carbides containing Nb and V, and obtaining a large amount of extremely fine precipitates with a grain size of less than 100 nm.
• the rolling conditions at 950°C or above, the cumulative rolling reduction rate at 900°C or below, the rolling conditions at 900°C or below, and the rolling end temperature are adjusted.
• intragranular dislocations are increased in both the rolling and accelerated cooling processes.
• the present invention ensures high strength in the mid-temperature range by increasing dislocations through rolling and accelerated cooling, and by suppressing dislocation recovery in the mid-temperature range through the fine carbides that are dispersed and precipitated by heating after accelerated cooling.
• the present invention has been completed based on the above findings. Specifically, the present invention provides the following. [1] In mass%, C: 0.040 to 0.090%, Si: 0.03-0.30%, Mn: 1.50-2.50%, P: 0.020% or less, S: 0.002% or less, Mo: 0.10-0.60%, Nb: 0.020-0.070%, Ti: 0.020% or less, V: 0.080% or less, Al: 0.045% or less, N: 0.010% or less, the balance being Fe and unavoidable impurities, and the parameter P eff represented by the following formula (1) is 0.050% or more,
• the microstructure has bainite in an area ratio of 80% or more at the center of the sheet thickness, the average grain size of the bainite is 30 ⁇ m or less, and the minimum grain size of the top 20% of the bainite grains having the largest grain size is 70 ⁇ m or less, A steel plate having a yield strength of 555 MPa or more before and after aging under the condition of Larson Miller Parameter
• the microstructure has an area ratio of ferrite of 10% or less and an area ratio of island martensite (MA) of 10% or less at the center of the plate thickness
• the steel sheet according to any one of [1] to [3], wherein at the sheet thickness center, 50,000 to 1,000,000 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 are present per mm2 of an observation area, and the difference in yield strength before and after aging, obtained by subtracting the yield strength after aging from the yield strength before aging, is 50
• [5] A steel pipe using the steel plate according to any one of [1] to [4].
• [6] The method for producing a steel sheet according to any one of [1] to [4], a heating step of heating the steel material to 1000 to 1200°C; a hot rolling step in which the steel material heated in the heating step is hot-rolled under conditions including one or more passes of rolling at a cumulative reduction of 50% or more at 900°C or less and a reduction per pass of 10% or more at 950°C or more, and one or more passes of rolling at a reduction per pass of 15% or more at 900°C or less, and an end temperature of the rolling being 850°C or less;
• the method for producing a steel plate includes: an accelerated cooling step in which the hot-rolled steel plate obtained in the hot-rolling step is accelerated-cooled under conditions of a cooling start temperature of 700°C or higher, an average cooling rate of 10°C/s or higher, and a cooling stop temperature of 250 to 550°C; and a reheating step in
• a method for manufacturing a steel pipe comprising: a cold forming step of cold-forming the steel plate according to any one of [1] to [4] into a tubular shape; a welding step of welding the butt joints where the ends of the steel plate formed into a tubular shape in the cold forming step are butted together; and a pipe expansion step of expanding the pipe.
• the present invention it is possible to obtain steel pipes that are large in diameter but have a yield strength of 555 MPa or more, a tensile strength of 620 MPa or more, and excellent low-temperature toughness after long-term holding in the medium temperature range, which are required for high-strength steel pipes for steam transportation. Furthermore, according to the present invention, it is possible to obtain steel pipes with the above properties even while reducing the amount of alloying elements used and keeping manufacturing costs down.
• the steel plate (also referred to as high-strength steel plate) of the present invention contains, in mass %, C: 0.040 to 0.090%, Si: 0.03 to 0.30%, Mn: 1.50 to 2.50%, P: 0.020% or less, S: 0.002% or less, Mo: 0.10 to 0.60%, Nb: 0.020 to 0.070%, Ti: 0.020% or less, V: 0.080% or less, Al: 0.045% or less, and N: 0.010% or less.
• “%" representing the content of a component means “mass %.”
• Carbon (C) is an element necessary for ensuring the strength of steel through solid solution strengthening and precipitation strengthening.
• an increase in the amount of solute C and the formation of precipitates contribute to ensuring strength in the intermediate temperature range.
• the present invention specifies a C content of 0.040% or more.
• the C content is preferably 0.045% or more, more preferably 0.048% or more, even more preferably 0.050% or more, and most preferably 0.053% or more.
• a C content exceeding 0.090% leads to deterioration in toughness and weldability. Therefore, the upper limit of the C content is set to 0.090%.
• the C content is preferably 0.080% or less, more preferably 0.075% or less, even more preferably 0.070% or less, and most preferably 0.065% or less.
• Si 0.03 ⁇ 0.30% Si is added for deoxidation. If the Si content is less than 0.03%, a sufficient deoxidation effect cannot be obtained. Therefore, the Si content is set to 0.03% or more.
• the Si content is preferably set to 0.04% or more, more preferably 0.05% or more, even more preferably 0.07% or more, and most preferably 0.10% or more.
• the Si content is set to 0.30% or less.
• the Si content is preferably set to 0.28% or less.
• the Si content is more preferably set to 0.25% or less, more preferably 0.23% or less, even more preferably 0.20% or less, and most preferably 0.18% or less.
• Mn 1.50-2.50%
• Mn is an element effective in improving the strength and toughness of steel. This effect can be fully achieved by setting the Mn content to 1.50% or more.
• the Mn content is preferably 1.55% or more.
• the Mn content is more preferably 1.60% or more, more preferably 1.65% or more, even more preferably 1.70% or more, and most preferably 1.75% or more.
• the Mn content is set to 2.50% or less.
• the Mn content is preferably 2.40% or less, more preferably 2.30% or less.
• the Mn content is more preferably 2.20% or less, and most preferably 2.10% or less.
• P 0.020% or less
• P is an impurity element that significantly 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 and can significantly deteriorate 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 morphology of S from MnS to CaS-based inclusions, in the case of high-strength steel plates of X80 grade or higher, finely dispersed CaS-based inclusions can also cause deterioration of toughness. Therefore, 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 0.001% 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 ⁇ 0.60% Mo contributes significantly to increasing strength at room temperature and in the intermediate temperature range by forming solid solutions or precipitates. However, if the Mo content is less than 0.10%, sufficient strength cannot be obtained in the intermediate temperature range, so 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 0.60%, toughness and weldability deteriorate, so the Mo content is set to 0.60% or less.
• the Mo content is preferably set to 0.55% or less, more preferably 0.50% or less, even more preferably 0.45% or less, and most preferably 0.40% or less.
• Nb 0.020-0.070%
• Nb is an important element in the present invention. Specifically, Nb is a component necessary for forming carbides and ensuring strength at room temperature and in the intermediate temperature range. Nb is also necessary for refining the microstructure and imparting sufficient strength and toughness by suppressing grain growth during slab heating and rolling. This effect is most pronounced when the Nb content is 0.020% or more, so the Nb content is set to 0.020% or more. The Nb content is preferably set to 0.023% or more, more preferably 0.025% or more, even more preferably 0.028% or more, and most preferably 0.030% or more.
• the Nb content is set to 0.070% or less.
• the Nb content is preferably set to 0.068% or less.
• the Nb content is more preferably 0.065% or less, further preferably 0.063% or less, and most preferably 0.060% or less.
• Ti 0.020% or less Ti forms TiN, which suppresses grain growth during slab heating and in weld heat-affected zones.
• Ti has the effect of refining the microstructure and improving toughness.
• 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. If the Ti content exceeds 0.020%, the presence of TiN makes it difficult to disperse and precipitate fine carbides, making it difficult to suppress strength degradation in the intermediate temperature range. Therefore, the Ti content is set to 0.020% or less. Furthermore, the Ti content is preferably 0.019% or less. The Ti content is more preferably 0.018% or less, even more preferably 0.017% or less, and most preferably 0.016% or less.
• V 0.080% or less V forms complex precipitates with Ti and Nb, contributing to increased strength. Furthermore, V-based carbides are less likely to aggregate and coarsen even when held at high temperatures for long periods of time, making V a useful element for ensuring high-temperature creep strength. If the desired high-temperature creep strength can be achieved by including elements other than V, the steel sheet of the present invention does not need to contain V.
• the lower limit of the V content is not particularly limited and may be 0%. To achieve this effect, the V content is preferably 0.005% or more, more preferably 0.008% or more, even more preferably 0.010% or more, and most preferably 0.013% or more.
• the V content is specified to be 0.080% or less. Furthermore, the V content is preferably 0.070% or less. The V content is more preferably 0.060% or less, further preferably 0.050% or less, and most preferably 0.040% or less.
• Al 0.045% or less Al is added as a deoxidizer.
• the Al content is preferably 0.015% or more, and more preferably 0.020% or more. If the Al content exceeds 0.045%, the cleanliness of the steel decreases and the toughness deteriorates. Therefore, the Al content is set to 0.045% or less. Furthermore, the Al content is preferably 0.043% or less. The Al content is more preferably 0.040% or less, even more preferably 0.038% or less, and most preferably 0.035% or less.
• N 0.010% or less N forms TiN together with Ti.
• TiN 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, improving the toughness of the weld heat-affected zone.
• the N content is preferably 0.002% or more, and more preferably 0.0025% or more.
• the N content is set to 0.010% or less, and preferably 0.006% or less.
• the N content is more preferably 0.0055% or less, and even more preferably 0.005% or less.
• P eff (%) 0.050% or more P eff is defined as (0.13Nb + 0.24V - 0.125Ti) / (C + 0.86N) ... (1).
• the element symbols represent the content (mass%) of each element, and 0 is substituted for elements that are not contained.
• P eff it is necessary to adjust the content of the above elements so that P eff is 0.050% or more.
• P eff is an important factor for ensuring that steel composed within the above composition range has excellent strength in the intermediate temperature range. If P eff (%) is less than 0.050%, the amount of finely dispersed carbides precipitated upon reheating after cooling will be small. As a result, strength, particularly tensile strength after long-term heat treatment, will decrease significantly.
• P eff (%) is set to 0.050% or more.
• P eff is preferably 0.055% or more to sufficiently suppress strength decrease after heat treatment.
• P eff (%) is more preferably 0.060% or more, even more preferably 0.065% or more, and most preferably 0.070% or more.
• P eff is preferably 0.280% or less, more preferably 0.270% or less, even more preferably 0.260% or less, and most preferably 0.250% or less.
• the steel sheet of the present invention may contain, in addition to the above-mentioned chemical composition, one or more of Cu, Ni, Cr, and Ca.
• Cu 0.50% or less
• the Cu content is preferably 0.05% or more, and more preferably 0.10% or more.
• a Cu content of more than 0.50% impairs weldability, so when Cu is contained, the Cu content is set to 0.50% or less.
• the Cu content is preferably 0.48% or less.
• the Cu content is 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.
• the Ni content is preferably 0.05% or more, and more preferably 0.10% or more. If the Ni content exceeds 0.50%, not only does the effect saturate, but it also leads to an increase in manufacturing costs. Therefore, when Ni is contained, its content is set to 0.50% or less.
• the Ni content is preferably set to 0.48% or less.
• the Ni content is more preferably set to 0.45% or less, even more preferably set to 0.43% or less, and most preferably set to 0.40% or less.
• Cr 0.50% or less Cr is one of the elements effective in increasing strength. To obtain this effect, the Cr content is preferably 0.05% or more, and more preferably 0.10% or more. A Cr content exceeding 0.50% has an adverse effect on 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.
• Ca controls the morphology of sulfide-based inclusions and improves toughness. This effect is manifested by setting the Ca content to 0.0005% or more. Therefore, when Ca is contained, the Ca content is set to 0.0005% or more.
• the Ca content is preferably set to 0.0010% or more. When the Ca content exceeds 0.0040%, not only does the above effect saturate, but cleanliness decreases and toughness deteriorates. Therefore, when Ca is contained, its 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.
• Ti/N 2.0 to 4.0 (preferable requirement)
• TiN is finely dispersed, and prior austenite grains in the weld heat affected zone are refined.
• the refinement of prior austenite grains in the weld heat affected zone improves the toughness of the weld heat affected zone in the low-temperature range of ⁇ 40°C or less and the medium-temperature range of 300°C or more.
• Ti/N ratio is less than 2.0, the effect is insufficient, so Ti/N is preferably 2.0 or more, more preferably 2.1 or more, even more preferably 2.2 or more, and most preferably 2.3 or more.
• Ti/N is preferably 4.0 or less, more preferably 3.9 or less, even more preferably 3.8 or less, and most preferably 3.7 or less.
• X 0.70% or more (preferable requirement)
• X 0.35Cr+0.9Mo+12Nb+8V...(3)
• the element symbols in formula (3) represent the content (mass%) of each element. Elements that are not contained are substituted with 0. Cr, Mo, Nb, and V contribute to improving temper softening resistance and intragranular precipitation strengthening during rolling in steels composed within the above-mentioned composition ranges.
• Formula (3) above is an important factor for obtaining steels with excellent strength of X80 grade or higher in the medium temperature range after long-term heat treatment and good low-temperature toughness. By combining this with the manufacturing conditions described below, the effect of satisfying formula (3) is greatly manifested.
• X is preferably 0.70% or more in the present invention. More preferably, X is 0.75% or more. To achieve X100 grade strength after long-term heat treatment at 350°C, X is even more preferably 0.90% or more. Most preferably, X is 1.00% or more. Furthermore, if X is 2.00% or more, the low-temperature toughness of the weld may decrease. Therefore, X is preferably less than 2.00%. Preferably, X is less than 1.90%, more preferably less than 1.80%, even more preferably 1.70% or less, and most preferably 1.60% or less.
• Y Cu+Ni+Cr+Mo...(4)
• the element symbols in formula (4) represent the content of each element, and 0 is substituted for elements that are not contained.
• Y in the above formula (4) is preferably 1.50% or less.
• the total content of the above elements is preferably 1.50% or less. It is more preferably 1.40% or less, even more preferably 1.30% or less, and most preferably 1.20% or less.
• the lower limit may be 0, but Y is preferably 0.20% or more, and more preferably 0.30% or more. Note that one of the features of the present invention is that desired properties can be obtained even when the amounts of these components used are reduced.
• 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.
• unavoidable impurities include Pb, Zn, Sn, As, B, Sb, Bi, Co, H, O, and REM.
• Bainite Area Fraction of 80% or More Bainite is an important structure for achieving both strength and low-temperature toughness. Bainite also effectively contributes to improving the strength of steel plates by strengthening the transformation structure. Furthermore, structural uniformity is necessary for reasons of increasing the initial dislocation density and from the viewpoint of improving the strength of high-strength steel plates, particularly strength in the mid-temperature range. Therefore, the structure of the steel plate of the present invention must be a structure predominantly composed of bainite. Specifically, the area fraction of bainite must be 80% or more of the entire steel structure at the center of the plate thickness.
• 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 the upper limit of the bainite fraction is not particularly limited, 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)
• 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.
• the area fraction of ferrite is preferably 1% or more, and more preferably 2% or more.
• Island martensite (MA: Martensite-Austenite constituent) is a very hard phase and may act as a fracture origin, 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 island martensite fraction, but an area fraction of 1% or more is preferred, and an area fraction of 2% or more is preferable.
• 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 30 ⁇ m or less, and the minimum grain size of the top 20% of the largest grains of the bainite is 70 ⁇ m or less. Because bainite grain boundaries act as resistance to brittle crack propagation, grain refinement contributes to improved low-temperature toughness. Therefore, the average grain size of bainite is 30 ⁇ m or less.
• the average grain size of bainite is preferably 28 ⁇ m or less.
• the average grain size of bainite is more preferably 25 ⁇ m or less, even more preferably 23 ⁇ m or less, and most preferably 20 ⁇ m or less. There is no particular lower limit, but it is preferably 5 ⁇ m or more, and more preferably 6 ⁇ m or more.
• the minimum grain size of the top 20% of grains with the largest grain size of bainite is preferably 60 ⁇ m or less, more preferably 55 ⁇ m or less, even more preferably 50 ⁇ m or less, and most preferably 45 ⁇ m or less. There is no particular lower limit, but 10 ⁇ m or more is preferred, and 15 ⁇ m or more is more preferred.
• the crystal orientation of a randomly selected 1 mm x 1 mm area at the center of the plate thickness was measured using electron backscatter diffraction (EBSD), and areas where the angle difference between adjacent pixels was 15° or more were determined to be grain boundaries through image analysis.
• EBSD electron backscatter diffraction
• the minimum grain size of the top 20% of bainite grains with the largest grain size represents the minimum grain size of the grains when the circle-equivalent diameter of each grain is used as the grain size and the largest 20% of the total number of grains is selected.
• Ti-based precipitates with a diameter of 100 nm or less there are 50,000 to 1,000,000 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 of Ti, Nb, V, Mo, Cr, and Al per mm2 of the test area (optimal conditions).
• 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
• the number of 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.
• 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 precipitates with a diameter of 100 nm or less 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.
• 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).
• SEM scanning electron microscope
• TEM transmission electron microscope
• the number of precipitates with a size of 100 nm or less is measured in an area of at least 100 ⁇ m2 . 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 ⁇ m2 , so 25 fields of view are observed randomly. This measurement result is converted to the number per unit area.
• 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).
• the tensile strength (TS) at 350°C measured after aging under the condition of Larson Miller Parameter (LMP) 15700 and the tensile strength ( TS0 ) at 350°C measured before aging satisfy the relationship ( TS0 - TS)/TS0 ⁇ 0.050.
• ( TS0 - TS)/ TS0 is an index for evaluating the decrease in tensile strength when held for a long time in the mid-temperature range. If this index is 0.050 or less, the decrease in tensile strength after long-term holding in the mid-temperature range falls within a range that is practically acceptable.
• (TS0 - TS)/ TS0 is set to 0.050 or less.
• (TS0 - TS)/ TS0 is preferably 0.049 or less. It is more preferably 0.048 or less, even more preferably 0.047 or less, and most preferably 0.046 or less.
• the lower limit is not particularly limited, and may be a negative value or may be ⁇ 0.100 or more.
• 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 condition).
• 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.
• 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, most preferably 35 MPa or less, and even more preferably 30 MPa or less.
• the lower limit is not particularly limited, and may be a negative value, or may be ⁇ 100 MPa or more.
• 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 value 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 aging treatment conditions under the above conditions include, for example, heat treatment at 400° C. for 2335 hours.
• the steel sheet of the present invention has a yield strength of 555 MPa or more and a tensile strength of 620 MPa or more, measured at 350°C.
• the yield strength and tensile strength measured at 350°C are preferably 560 MPa or more and 625 MPa or more, respectively, more preferably 565 MPa or more and 630 MPa or more, respectively, and even more preferably 570 MPa or more and 635 MPa or more, respectively. While there are no particular upper limits, the yield strength and tensile strength measured at 350°C are preferably 840 MPa or less and 900 MPa or less, respectively, and more preferably 830 MPa or less and 890 MPa or less, respectively.
• the yield strength and tensile strength measured at 350°C after long-term aging in the medium temperature range are 555 MPa or more and 620 MPa or more.
• the yield strength and tensile strength measured at 350°C after long-term aging in the mid-temperature range are preferably 560 MPa or more and 625 MPa or more, respectively, more preferably 565 MPa or more and 630 MPa or more, respectively, and even more preferably 570 MPa or more and 635 MPa or more, respectively.
• the yield strength and tensile strength measured at 350°C after long-term aging in the mid-temperature range are preferably 840 MPa or less and 900 MPa or less, respectively, and more preferably 790 MPa or less and 850 MPa or less, respectively. These excellent physical properties can be achieved by adjusting the specific component composition and employing the manufacturing conditions described below.
• Toughness of Steel Plate Ductile Fracture Area Ratio Obtained by DWTT at -40°C of 85% or More
• the toughness of the steel plate of the present invention is such that the ductile fracture area ratio (DWTTSA-40°C) obtained by DWTT (test temperature: -40°C) in accordance with API 5L is 85% or more. If this value is less than 85%, the steel plate will undergo brittle fracture at low temperatures, making it difficult to use in year-round installation work, including winter when the temperature drops below 0°C, or in areas with extremely low ambient temperatures. This value must be achieved because it is difficult to use. This also means that the fracture transition temperature is -40°C or less.
• the test temperature in the DWTT is set to -40°C to take into account the decrease in toughness due to work hardening during pipe making. Furthermore, the ductile fracture area ratio is preferably 86% or more, more preferably 87% or more, and even more preferably 88% or more. The upper limit is not particularly limited, and the ductile fracture area ratio obtained by DWTT at -40°C may be 100% or less.
• Toughness of weld heat affected zone is 60J or more.
• the toughness of the weld heat affected zone (HAZ) formed when the steel plate of the present invention is welded to the same steel plate or to another steel is 60J or more in terms of absorbed energy vE -40 when performed in a Charpy impact test at a test temperature of -40°C. If vE - 40 is 60J or more, the toughness required for structural pipe can be ensured.
• the absorbed energy vE-40 is preferably 70J or more, more preferably 80J or more, and even more preferably 100J or more. There is no particular upper limit, but it is preferably 400J or less, more preferably 380J or less, and even more preferably 350J or less.
• the notch position of the Charpy impact test specimen is 3mm toward the base metal from the bond line, which is the boundary between the weld metal and the base metal (HAZ 3mm). Furthermore, when a Charpy impact test is performed using three test pieces for each condition, the average value of absorbed energy (vE -40 ) is 60 J or more, which is considered to be within the scope of the present invention.
• the upper limit is not particularly limited.
• the welding method is submerged arc welding, and the welding heat input is 85 kJ/cm or less, depending on the thickness of the steel plate.
• the heat input is preferably 80 kJ/cm or less, and more preferably 70 kJ/cm or less.
• the lower limit is not particularly limited, but is preferably 20 kJ/cm or more, and more preferably 25 kJ/cm or more.
• the steel pipe of the present invention is manufactured using the 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 transporting steam, even when it has a large diameter.
• "Large diameter” means that the outer diameter (diameter) of the steel pipe (also called high-strength 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 of high-strength welded steel pipe for steam transportation.
• the thickness of the steel pipe is not particularly limited, but for 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 method for manufacturing a steel plate according to the present invention includes a heating step, a hot rolling step, an accelerated cooling step, and a reheating step.
• the temperatures in the description of each step refer to the average temperature in the thickness direction of the steel plate unless otherwise specified.
• the average temperature in the thickness direction is determined by simulation calculations using parameters such as the plate thickness, cooling conditions, and heat transfer coefficient from the surface temperature of the slab (steel material) or the steel plate.
• 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 cooling rate is the average 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 end of cooling.
• the heating rate (heating rate) in the reheating step is the average heating rate obtained by dividing the temperature difference required for reheating to the reheating temperature after cooling in the accelerated cooling step by the time required for reheating.
• the heating process is a process of heating a steel material to 1000 to 1200°C.
• the steel material is, for example, a slab obtained by casting molten steel. Since the composition of the steel material determines the composition of the steel plate, the composition of the steel plate can be adjusted at the stage of adjusting the composition of the molten steel.
• the steelmaking method for the steel material is not particularly limited. In the hot rolling process described below, the heating temperature is set to 1000°C or higher to sufficiently promote austenitization and carbide solid solution and obtain sufficient strength at room temperature and in the intermediate temperature range.
• 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.
• the heating temperature is set to 1200°C or lower.
• the heating temperature is preferably 1190°C or less, more preferably 1180°C or less, even more preferably 1170°C or less, and most preferably 1160°C or less.
• the hot rolling step is a step in which the steel material heated in the heating step is hot rolled under conditions including one or more passes of rolling at a cumulative reduction rate of 50% or more at 900°C or less and a reduction rate per pass of 10% or more at 950°C or more, and one or more passes of rolling at 900°C or less and a reduction rate per pass of 15% or more, and the rolling finish temperature is 850°C or less.
• the hot rolling step is an important manufacturing condition of the present invention.
• austenite grains are elongated and become finer in the thickness and width directions, and the density of dislocations within the grains introduced by rolling is increased.
• This effect is achieved by including at least one pass of rolling at a cumulative reduction of 50% or more at 900°C or less, at least one pass of rolling at a reduction per pass of 10% or more at 950°C or more, at least one pass of rolling at a reduction per pass of 15% or more at 900°C or less, and by setting the rolling end temperature to 850°C or less.
• the cumulative reduction rate at 900°C or below is set to 50% or more.
• a cumulative reduction rate of 55% or more is preferable, a cumulative reduction rate of 60% or more is more preferable, a cumulative reduction rate of 65% or more is even more preferable, and a cumulative reduction rate of 70% or more is most preferable.
• the rolling end temperature is set to 850°C or lower.
• the rolling end temperature is preferably 840°C or lower, more preferably 830°C or lower, even more preferably 820°C or lower, and most preferably 810°C or lower.
• the cumulative reduction is preferably 95% or less, more preferably 90% or less, and even more preferably 85% or less.
• the rolling end temperature is preferably Ar3 temperature or higher.
• [C], [Mn], [Cu], [Ni], [Cr] and [Mo] respectively represent the contents (mass%) of C, Mn, Cu, Ni, Cr and Mo in the steel of the base steel plate.
• the content of the element may be set to "0" to determine the Ar3 temperature.
• austenite grains are refined through recrystallization, thereby improving the low-temperature toughness of the steel sheet. 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.
• the number of passes for rolling at a reduction rate of 10% or more per pass may be 50 passes or less, more preferably 30 passes or less, and even more preferably 15 passes or less.
• the reduction rate per pass at 950°C or higher, but it is preferably 40% or less.
• the number of passes of rolling at a reduction rate of 15% or more per pass there is no particular upper limit on the number of passes of 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. It is also preferable to include one or more passes of rolling at 900°C or below with a reduction rate of 20% or more per pass. Furthermore, there is no upper limit to the rolling reduction per pass at temperatures below 900°C, but it is preferable that it be 40% or less.
• the accelerated cooling step is a step of accelerated cooling the hot-rolled steel sheet obtained in the hot rolling step under the conditions of a cooling start temperature of 700°C or higher, an average cooling rate of 10°C/s or higher, and a cooling stop temperature of 250 to 550°C.
• the average cooling rate means a 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.
• the cooling start temperature 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 must be 700°C or higher.
• the cooling start temperature is preferably 710°C or higher, more preferably 720°C or higher, even more preferably 730°C or higher, and most preferably 740°C or higher. There is no particular upper limit, but it is preferably 850°C or lower, and more preferably 800°C or lower.
• the strength of 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 10°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. For this reason, if the average cooling rate is less than 10°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 10°C/s, the effect of refining the structure is reduced, the crystal grain size does not decrease, and low-temperature toughness deteriorates. For this reason, the average cooling rate during accelerated cooling is set to 10°C/s or more.
• the average cooling rate is preferably 12°C/s or more, more preferably 14°C/s or more, even more preferably 16°C/s or more, and most preferably 18°C/s or more. There is no particular upper limit to the average cooling rate, but 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.
• the strength of steel plate 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. As a result, sufficient strength, particularly in the mid-temperature range, cannot be obtained. Therefore, the cooling stop temperature is set to 550°C or lower.
• the cooling stop temperature is preferably set to 540°C or lower, more preferably 530°C or lower, even more preferably 520°C or lower, and most preferably 450°C or lower.
• the cooling stop temperature of accelerated cooling is set to 250°C or higher.
• the cooling stop temperature is preferably set to 260°C or higher, more preferably 270°C or higher, even more preferably 280°C or higher, and most preferably 290°C or higher.
• the reheating process is a process in which the hot-rolled steel sheet is reheated within 150 seconds (s) after cooling is stopped in the accelerated cooling process, under conditions of a heating rate of 1°C/s or more and an ultimate temperature of 550 to 700°C.
• the time from the stop of cooling in the accelerated cooling process to the start of the reheating process is 150 seconds or less, preferably 130 seconds or less, more preferably 120 seconds or less, even more preferably 110 seconds or less, and most preferably 100 seconds or less.
• the lower limit of the time from the stop of cooling in the accelerated cooling process to the start of the reheating process is not particularly limited, but is preferably 5 seconds or more, more preferably 6 seconds or more, and even more preferably 7 seconds or more.
• the reheating process after the accelerated cooling process in which the heating rate is set to 1°C/s or more and the ultimate temperature is set to 550-700°C, is important in this invention.
• This process allows fine precipitates, which contribute to strengthening at room temperature and in the intermediate temperature range, to precipitate during reheating.
• the material must be reheated to a temperature range of 550-700°C immediately after accelerated cooling. There is no need to set a specific temperature holding time during the reheating process.
• air cooling including air blast cooling is used after reheating.
• the heating rate in the reheating step after accelerated cooling should be 1°C/s or more, preferably 3°C/s or more, more preferably 5°C/s or more, even more preferably 6°C/s or more, and most preferably 7°C/s or more.
• the heating rate in the reheating step is preferably 100°C/s or less, more preferably 75°C/s or less, and even more preferably 50°C/s or less.
• the reheating temperature is set to 550°C or higher.
• the reheating temperature is preferably 560°C or higher, more preferably 570°C or higher, even more preferably 580°C or higher, and most preferably 590°C or higher.
• the reheating temperature exceeds 700°C, the precipitates become coarse and sufficient strength cannot be obtained at room temperature or in the mid-temperature range, so the reheating temperature is set to 700°C or lower.
• the reheating temperature is preferably 690°C or lower, more preferably 680°C or lower, even more preferably 670°C or lower, and most preferably 660°C or lower. It should be noted that the temperature increase rate of 1°C/s or more in the reheating step after accelerated cooling specified in the present invention is difficult to achieve in an atmospheric furnace depending on the plate thickness. Therefore, it is preferable to use a gas combustion furnace or induction heating device that is capable of rapidly heating the steel plate as the heating device. It is more preferable to install the gas combustion furnace or induction heating device on the conveying line downstream of the cooling equipment for performing accelerated cooling.
• Induction heating devices have easier temperature control than soaking furnaces and are relatively inexpensive. Furthermore, induction heating devices are particularly preferable because they can quickly heat steel sheets after they have been cooled. Furthermore, by arranging multiple induction heating devices in series, it is possible to freely control the heating rate and reheating temperature, even when the line speed or type or size of steel sheets varies, simply by setting the number of induction heating devices and the power supply.
• the method for producing a steel pipe of the present invention includes a cold forming step, a welding step, and a pipe expansion step.
• the cold forming process is a process of cold forming the steel plate of the present invention into a tubular shape.
• the thickness of the steel plate is preferably 12 mm or more. Also, it is preferably 30 mm or less. More preferred ranges are as described above.
• the method for cold forming the steel plate into a tubular shape is not particularly limited. Examples of forming methods include methods for forming into a steel pipe shape by cold forming such as the UOE process, press bending (also called bending press), and roll forming.
• grooves are formed on the widthwise ends of the raw steel plate, and then a press is used to bend the widthwise ends of the steel plate.
• the steel plate is then formed into a U-shape and then an O-shape using the press, thereby forming the steel plate into a cylindrical shape so that the widthwise ends of the steel plate face each other.
• press bending a steel plate is successively formed by repeatedly three-point bending to produce a steel pipe having a substantially circular cross section.
• the welding process is a process of welding the butt joints of steel plates formed into a tubular shape in the cold forming process.
• the welding method is not particularly limited, but submerged arc welding or the like may be used. Opposing widthwise ends of the steel plates are butt-welded. This welding is called seam welding.
• a two-step process is preferred: a tack welding process in which the cylindrical steel plates are restrained and the widthwise ends of the steel plates are butt-welded together, and a main welding process in which welding is performed on the inner and outer surfaces of the butt joints of the steel plates using submerged arc welding.
• the pipe expansion ratio (the change in outer diameter of the steel pipe before and after expansion divided by the outer diameter of the steel pipe before expansion) is usually preferably 0.3% or more. Also, the pipe expansion ratio is preferably 1.5% or less. Furthermore, from the viewpoint of balancing the roundness improvement effect and the capacity required of the pipe expansion device, the pipe expansion ratio is more preferably 0.5% or more. And the pipe expansion ratio is more preferably 1.2% or less.
• Heat treatment after steel pipe production can be carried out according to the desired properties, and there are no specific regulations.
• a 1.0% expansion ratio means that the inner diameter of the pipe was expanded by 1.0% in the radial direction from the inner surface to the outer surface using a pipe expander. Furthermore, a reheating process was performed within 150 seconds (s) after cooling was stopped in the accelerated cooling process at the heating rate shown in Tables 2-1 and 2-2.
• a sample for observing the steel structure was taken from the center of the width of the steel plate manufactured as described above, and the L-section of the steel plate (a cross section parallel to the rolling direction and normal to the rolling surface) was mirror-polished.
• the crystal orientation of a randomly selected 1 mm x 1 mm area at the center of the plate thickness was then measured using electron backscatter diffraction (EBSD).A region where the angle difference between adjacent pixels was 15° or more was considered to be a grain boundary, and the grain size was determined by image analysis.
• the EBSD measurement conditions were an acceleration voltage of 17 kV and a measurement pitch of 0.8 ⁇ m.
• the minimum grain size of the top 20% of bainite grains in terms of grain size represents the minimum grain size of the grains when the circle-equivalent diameter of each grain is used as the grain size and the largest 20% of the total number of grains is selected.
• 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, steel structure photographs were taken of five randomly selected 7.1 ⁇ 10 ⁇ 2 mm 2 regions (magnification: 400x) at the center of the sheet thickness.
• the bainite fraction, ferrite fraction, and island martensite fraction in the photographs were measured using an image analyzer (Fiji), and the bainite grain size was measured using an image analyzer (OIM Analysis, manufactured by TSL).
• the bainite fraction, ferrite fraction, island martensite fraction, and bainite grain size were averaged over the five fields of view. Note that in the photographs, regions that were not elongated in the rolling direction and were observed as equiaxed crystal grains were determined to be bainite.
• the number of precipitates was determined using the evaluation method used in the embodiment.
• DWTT test specimens were taken so that the longitudinal direction of the DWTT test specimen was horizontal to the rolling direction of the steel plate and perpendicular to the plate thickness direction, and DWTT in accordance with API 5L was performed.
• the ductile fracture surface area ratio and fracture transition temperature at a test temperature of -40°C were evaluated.
• a ductile fracture surface area ratio of 85% or more and a fracture transition temperature of -40°C or less at a test temperature of -40°C were evaluated as good.
• tensile test pieces were taken so that the longitudinal direction of the tensile test piece 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 and tensile strength.
• round bar test pieces with a diameter of 6 mm were used, and the crosshead speed was 0.15 mm/min.
• a yield strength of 555 MPa or more and a tensile strength of 620 MPa or more at 350°C were evaluated as good. Note that the steel sheet properties were evaluated by taking test pieces from the steel sheet before it was formed into a steel pipe.
• the yield strength and tensile strength at 350°C were measured after heat treatment (aging treatment) in a N (nitrogen) atmosphere furnace under conditions of Larson-Miller Parameter 15700 (400°C, 2335 hours), which is the tempering parameter shown in equation (2), equivalent to 20 years of retention at 350°C, the applicable temperature for steam piping.
• Steel plates were evaluated as good if their yield strength at 350°C after heat treatment was 555 MPa or more and their tensile strength was 620 MPa or more.
• the above measurements were performed on both the steel plate and steel pipe in the same manner as before heat treatment. The results are shown in Tables 3-1 and 3-2.
• the tensile strength of the steel pipe properties was calculated as ((tensile strength before heat treatment (TS 0 )) - (tensile strength after heat treatment (TS)))/tensile strength before heat treatment (TS 0 ), and a value of 0.050 or less was evaluated as good.
• the toughness of the weld heat-affected zone was evaluated by a Charpy impact test.
• the notch position of the Charpy impact test specimen was 3 mm toward the base metal from the bond line, which is the boundary between the weld metal and the base metal (HAZ 3 mm).
• the test was carried out at a temperature of -40°C.
• the Charpy impact test was carried out using three specimens for each condition, and specimens with an average absorbed energy (vE -40 ) of 60 J or more at -40°C were evaluated as having excellent toughness.
• Tables 2-1, 2-2, 3-1, and 3-2 show the manufacturing conditions for the steel plates and steel pipes, as well as the test results.
• the inventive examples (Nos. 1 to 34) the low-temperature toughness of the steel plate base material at -40°C, the HAZ toughness, and the ( TS0 -TS)/ TS0 ratio were all good.
• the yield strength, tensile strength, low-temperature toughness at -40°C, HAZ toughness, and ( TS0 -TS)/ TS0 before and after the heat treatment of the steel plate did not reach the target value.

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