Steel sheet, steel tube, method for producing steel sheet, and method for producing steel tube

Abstract

The present invention provides a steel sheet serving as a material of a steel pipe, and a method for producing the steel sheet.　A steel sheet according to the present invention has a specific composition and a specific structure. The toughness of the steel sheet is 85% or greater at a ductile fracture rate obtained by DWTT at −40°C. The toughness of a weld-heat-affected zone formed during welding is 60 J or greater in vE-40. The yield strength before and after aging performed under a condition in which Larson Miller Parameter (LMP) = 15700, which is defined by formula (2), is 555 MPa or greater, and the tensile strength (TS) at 350°C measured after aging and the tensile strength (TS0) at 350°C measured before aging satisfy the relationship (TS0 − TS) / TS0 ≤ 0.050. Formula 2: LMP = (T + 273) × (20 + log(t)) (T: heat treatment temperature (°C); t: heat treatment time (hours))

Description

Steel plate, steel pipe, method for manufacturing steel plate and method for manufacturing steel pipe

[0001] 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 at medium temperatures, a method for manufacturing the same, and a steel pipe made from the steel plate and a method for manufacturing the same. In particular, the 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 disclose methods for manufacturing steel pipes for steam transportation that can be used in the steam injection method. These patent documents disclose seamless steel pipes equivalent to API X80 grade, but the outer diameter of these seamless steel pipes is a maximum of 16 inches, and it is difficult to further increase the diameter of the seamless steel pipes.

[0004] Furthermore, in seamless steel pipes, the addition of large amounts of alloying elements is necessary to obtain strength of API X80 grade or higher. In recent years, with regard to manufacturing techniques for high-strength steel pipes that are manufactured by welding and can be made larger in diameter, manufacturing techniques for high-strength steel pipes having strength of API X80 or higher have been disclosed in Patent Documents 3, 4, and 5.

[0005] Japanese Patent Publication No. 2000-290728 Japanese Patent No. 4821939 Japanese Patent No. 5055736 International Publication No. 2012 / 108027 Japanese Patent No. 6137435

[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 discloses a manufacturing technique for high-strength steel plate having a yield strength of 700 MPa or more. However, in order to ensure strength in the intermediate temperature range using the manufacturing technique for high-strength steel plate disclosed in Patent Document 4, it is necessary to add a large amount of alloying elements. Furthermore, it became clear during the process of completing the present invention that the tensile strength of the high-strength steel plate described in Patent Document 4 is significantly reduced when held in the intermediate temperature range for a long period of time.

[0008] The high-strength steel sheet described in Patent Document 5 satisfies the API X80 grade in terms of high-temperature properties in the medium temperature range, and the tensile strength (TS) at 350°C measured after aging under the condition of LMP = 15700 is higher than the tensile strength (TS) at 350°C measured before aging. ０ ) is (TS ０ -TS) / TS ０ Although the relationship of ≦0.050 is satisfied, the structure of the steel plate is not specified, and low temperature toughness is not taken into consideration.

[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] Steel pipes of API X80 or higher are required to satisfy 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 properties required for high-strength steel pipes for steam transport include satisfying the above strength properties 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 weld heat-affected zone formed during welding.

[0011] The present invention has been made to solve the above-mentioned problems, and an object of the present invention is to provide a steel plate that serves as a material for the steel pipe, and a manufacturing method thereof. The steel plate must satisfy all of the following requirements: a yield strength of 555 MPa or more and a tensile strength of 620 MPa or more both before and after aging, 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. Another object of the present invention is to provide a steel pipe made from the above-mentioned steel plate, and a manufacturing method thereof.

[0012] 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.

[0013] In addition, "excellent toughness of the weld heat affected zone formed during welding" refers to the absorbed energy (vE) at -40°C in a Charpy impact test using a test piece taken from the weld heat affected zone. －４０ ) is 60J or more.

[0014] Furthermore, "large diameter" means that the outer diameter of the steel pipe is 400 mm or more.

[0015] The present inventors have conducted extensive research into the properties of steel plates for large-diameter steel pipes in the intermediate temperature range. As a result, they have found that by appropriately selecting the chemical composition and manufacturing conditions, a steel plate can be obtained that can be used to manufacture steel pipes having the strength characteristics and low-temperature toughness required for high-strength steel pipes for steam transportation, even though they are large in diameter. They have discovered that reheating during the bainite transformation of Nb-based steel containing Nb as a solid solution or Nb-V-based steel containing Nb and V as a solid solution in the manufacturing process of accelerated cooling after hot rolling and subsequent reheating not only strengthens the steel through bainite transformation during accelerated cooling, but also strengthens the steel through precipitation strengthening due to fine precipitates precipitated from bainite and untransformed austenite during reheating, and suppresses dislocation recovery in the intermediate temperature range, thereby suppressing strength degradation in the intermediate temperature range. Furthermore, the presence of TiN makes it difficult for Nb to dissolve. As a result, compared to when Ti is not added, fine Nb carbides are less likely to disperse and precipitate during reheating after accelerated cooling, making it more difficult to suppress strength degradation in the intermediate temperature range. However, P calculated by the following formula (1) ｅｆｆ When the P value is 0.050% or more, fine Nb carbides and V carbides are dispersed and precipitated sufficiently during reheating even when Ti is added, making it possible to suppress a decrease in strength in the medium temperature range. Nb carbides and V carbides refer to simple carbides containing Nb or V, and composite carbides containing Nb and V. 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.

[0016] Nb and V are elements that form carbides in steel. Strengthening steel by the precipitation of NbC has been a common practice. Furthermore, V-based carbides are 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. In the present invention, 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 intermediate temperature range. In the present invention, during reheating after accelerated cooling, the hot-rolled steel sheet is heated in an atmospheric furnace at a heating rate faster than the conventional heating rate used industrially. This suppresses the growth of Nb-containing carbides or carbides containing Nb and V, and produces a large amount of extremely fine precipitates with a grain size of less than 100 nm.

[0017] Furthermore, when producing the steel sheet of the present invention, in order to introduce a large amount of dislocations into the intragranular structure, prior to the dispersed precipitation of fine carbides by reheating after the accelerated cooling step, the rolling conditions at 950°C or higher, the cumulative rolling reduction rate at 900°C or lower, the rolling conditions at 900°C or lower, and the rolling end temperature are adjusted. In other words, when producing the steel sheet of the present invention, intragranular dislocations are increased in both the rolling and accelerated cooling steps.

[0018] As described above, the present invention ensures high strength in the intermediate temperature range by increasing dislocations through rolling and accelerated cooling, and by suppressing the recovery of dislocations in the intermediate temperature range through the fine carbides that are dispersed and precipitated by heating after accelerated cooling.

[0019] The present invention has been completed based on the above findings. Specifically, the present invention provides the following: [1] A steel sheet having a component composition containing, by 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, N: 0.010% or less, with the balance being Fe and unavoidable impurities, and having a parameter P represented by the following formula (1): ｅｆｆThe 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 μm or less, and the minimum grain size of the top 20% of the largest grains of the bainite is 70 μm or less, the yield strength before and after aging under the condition of Larson Miller Parameter (LMP) = 15700 defined by the following formula (2) is 555 MPa or more, and the difference between the tensile strength (TS) at 350 ° C. measured after aging and the tensile strength (TS) at 350 ° C. measured before the aging is 0.050% or more, ０ ) is (TS ０ -TS) / TS ０ ≦0.050, the toughness of the steel plate is 85% or more in terms of the ductile fracture area ratio obtained by DWTT at -40 ° C, and the toughness of the weld heat affected zone formed during welding is vE －４０ Steel plate with a strength of 60J or more. ｅｆｆ(%) = (0.13Nb + 0.24V - 0.125Ti) / (C + 0.86N) (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 steel sheet according to [1], wherein Ti / N is 2.0 to 4.0, and X represented by formula (3) is 0.70% or more. X = 0.35Cr + 0.9Mo + 12Nb + 8V (3) The element symbols in formula (3) represent the content (mass%) of each element. Also, 0 is substituted for elements that are not contained. [3] The steel sheet according to [1] or [2], wherein the composition is in mass %, and further contains one or more of Cu: 0.50% or less, Ni: 0.50% or less, Cr: 0.50% or less, and Ca: 0.0005 to 0.0040%, and Y represented by formula (4) is 1.50% or less. Y = Cu + Ni + Cr + Mo ... (4) The element symbols in formula (4) represent the content (mass %) of each element. Furthermore, 0 is substituted for elements that are not contained. [4] 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 sheet thickness, and 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 in diameter containing two or more elements of Ti, Nb, V, Mo, Cr, and Al are present in an inspection area of ​​1 mm at the center of the sheet thickness. ２[5] A steel pipe using the steel plate according to any of [1] to [4]. [6] A method for producing the steel plate according to any of [1] to [4], comprising: a heating step of heating a steel material to 1000 to 1200°C; and a hot rolling step of hot-rolling the steel material heated in the heating step under conditions of 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 a reduction per pass of 15% or more at 900°C or less, the hot-rolling finish temperature being 850°C or less. a reheating step of reheating the hot-rolled steel sheet obtained in the hot rolling step at a temperature of 550 to 700°C within 150 seconds after the cooling is stopped in the accelerated cooling step, at a temperature of 1°C / s or more and an ultimate temperature of 550 to 700°C, to obtain a steel sheet. [7] A method of manufacturing a steel pipe, comprising: a cold-forming step of cold-forming the steel sheet according to any one of [1] to [4] above into a tubular shape, a welding step of butting together the ends of the steel sheet formed into a tubular shape in the cold-forming step, and a pipe expanding step of expanding the pipe.

[0020] According to the present invention, it is possible to obtain a steel pipe having a large diameter and having 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 intermediate temperature range, which are required for high-strength steel pipes for steam transportation. Furthermore, according to the present invention, it is possible to obtain a steel pipe having the above properties even when the amount of alloying elements used is reduced and manufacturing costs are kept low.

[0021] Hereinafter, embodiments of the present invention will be described, but the present invention is not limited to the following embodiments.

[0022] <Steel Plate> 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. In the following description, "%" representing the content of a component means "mass%."

[0023] C: 0.040 to 0.090% 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 0.040% 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. Furthermore, a C content exceeding 0.090% leads to deterioration in toughness and weldability. For this reason, 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.

[0024] Si: 0.03 to 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. On the other hand, if the Si content exceeds 0.30%, toughness will deteriorate. Therefore, the Si content is set to 0.30% or less. Furthermore, 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.

[0025] Mn: 1.50 to 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. Furthermore, if the Mn content exceeds 2.50%, toughness and weldability will deteriorate significantly. Therefore, 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.

[0026] 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.

[0027] 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.

[0028] Mo: 0.10 to 0.60% Mo contributes significantly to increasing strength at room temperature and in the intermediate temperature range by forming solid solutions or precipitates. However, a Mo content of less than 0.10% does not provide sufficient strength in the intermediate temperature range, so 0.10% or more is included. Furthermore, the Mo content is preferably 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, a Mo content exceeding 0.60% degrades toughness and weldability, so it is set to 0.60% or less. Furthermore, the Mo content is preferably 0.55% or less, more preferably 0.50% or less, even more preferably 0.45% or less, and most preferably 0.40% or less.

[0029] Nb: 0.020 to 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 0.023% or more, more preferably 0.025% or more, even more preferably 0.028% or more, and most preferably 0.030% or more. If the Nb content exceeds 0.070%, not only does the effect nearly saturate but toughness also deteriorates, so the Nb content is set to 0.070% or less. The Nb content is preferably 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.

[0030] Ti: 0.020% or less Ti forms TiN, which suppresses grain growth during slab heating and in weld heat-affected zones. Thus, Ti has the effect of refining the microstructure and improving toughness. To achieve this effect, 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.

[0031] 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. On the other hand, if the V content exceeds 0.080%, the toughness of the weld heat-affected zone deteriorates. Therefore, 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.

[0032] Al: 0.045% or less Al is added as a deoxidizer. To obtain the effect of the 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.

[0033] 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. To achieve this effect, the N content is preferably 0.002% or more, and more preferably 0.0025% or more. Furthermore, if the N content exceeds 0.010%, the toughness of the base material deteriorates due to coarsening of precipitates and an increase in solute N, and further, the toughness of the weld metal when made into a steel pipe deteriorates. Therefore, 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.

[0034] P ｅｆｆ (%): 0.050% or more P ｅｆｆ is defined as (0.13Nb+0.24V-0.125Ti) / (C+0.86N) (1). In this formula (1), the element symbols represent the content (mass%) of each element, and 0 is substituted for elements that are not contained. P ｅｆｆ In the present invention, it is necessary to adjust the contents of the above elements so that P is 0.050% or more. ｅｆｆ is an important factor for making the steel composed of the above-mentioned composition range into a steel having excellent strength in the medium temperature range. ｅｆｆ If P (%) is less than 0.050%, the amount of finely dispersed carbides that precipitate when reheated after cooling is reduced. As a result, the strength, particularly the tensile strength after a long-term heat treatment, is significantly reduced. ｅｆｆ(%) is 0.050% or more. ｅｆｆ In order to sufficiently suppress the decrease in strength after heat treatment, P is preferably 0.055% or more. ｅｆｆ (%) is more preferably 0.060% or more, even more preferably 0.065% or more, and most preferably 0.070% or more. In addition, in order to prevent a large amount of precipitates from being generated in the weld heat affected zone, which would deteriorate the toughness, P is 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.

[0035] In order to further improve the properties, the steel sheet of the present invention may further contain, in addition to the above-mentioned composition, one or more of Cu, Ni, Cr and Ca.

[0036] Cu: 0.50% or less Cu is one of the elements effective in improving toughness and increasing strength. To achieve this effect, the Cu content is preferably 0.05% or more, and more preferably 0.10% or more. On the other hand, 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.

[0037] Ni: 0.50% or less Ni is one of the elements effective in improving toughness and increasing strength. To obtain this effect, 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.

[0038] 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.

[0039] Ca: 0.0005 to 0.0040% 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, if Ca is contained, the Ca content should be 0.0005% or more. The Ca content is preferably 0.0010% or more. If the Ca content exceeds 0.0040%, not only will the above effect saturate, but cleanliness will decrease and toughness will deteriorate. Therefore, if Ca is contained, its content should be 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.

[0040] Ti / N: 2.0 to 4.0 (preferred requirement) By specifying the Ti / N ratio within an appropriate range, TiN is finely dispersed, achieving refinement of prior austenite grains in the weld heat affected zone. Refining the 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. If the Ti / N ratio is less than 2.0, the effect is insufficient. Therefore, 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. If the Ti / N ratio exceeds 4.0, coarsening of the prior austenite grains due to coarsening of precipitates occurs. Since the coarsening of the prior austenite grains deteriorates the toughness of the weld heat affected zone, 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.

[0041] X: 0.70% or more (preferred requirement) X = 0.35Cr + 0.9Mo + 12Nb + 8V (3) The element symbols in formula (3) represent the content (mass%) of each element. Also, 0 is substituted for elements that are not contained. Cr, Mo, Nb, and V improve temper softening resistance and contribute to intragranular precipitation strengthening during rolling for steels composed within the above 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. The effect of satisfying formula (3) is greatly manifested by combining it with the manufacturing conditions described below. To achieve X80 grade strength after long-term heat treatment at 350°C, 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, it is more preferable that X be 0.90% or more. Most preferably, it 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.

[0042] Y: 1.50% or less (preferred requirement) Y = Cu + Ni + Cr + Mo (4) The element symbols in formula (4) represent the content of each element, with 0 being substituted for elements that are not contained. Y in formula (4) is preferably 1.50% or less. In order to achieve both the strength-increasing effect of these elements and production costs, 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. While the lower limit may be 0, Y is preferably 0.20% or more, and more preferably 0.30% or more. One of the features of the present invention is that desired properties can be obtained even when the amounts of these components are reduced.

[0043] 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.

[0044] Next, the structure of the steel sheet of the present invention will be described.

[0045] 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. To ensure a balance between strength and toughness, 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.

[0046] 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, and more preferably 2% or more.

[0047] Island martensite (MA): area fraction of 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 island martensite fraction, but an area fraction of 1% or more is preferred, and an area fraction of 2% or more is preferable. 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.

[0048] The average grain size of bainite is 30 μm or less, and the minimum grain size of the top 20% of the largest grains of the bainite is 70 μm 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 μm or less. The average grain size of bainite is preferably 28 μm or less. The average grain size of bainite is more preferably 25 μm or less, even more preferably 23 μm or less, and most preferably 20 μm or less. There is no particular lower limit, but it is preferably 5 μm or more, and more preferably 6 μm or more.

[0049] Furthermore, while reducing the average grain size improves low-temperature toughness, there is a limit to how much the average grain size can be reduced when considering factors such as plate crown or when cooling is initiated from a high temperature. Furthermore, in the present invention, it is essential to suppress the formation of coarse grains. Coarse bainite is likely to become the initiation point of fracture, and if the minimum grain size of the top 20% of grains with the largest grain size is large, low-temperature toughness deteriorates. In particular, when the minimum grain size of the top 20% of grains with the largest grain size exceeds 70 μm, they are likely to become the initiation point of fracture. Therefore, it is necessary to set the minimum grain size of the top 20% of grains with the largest grain size at the center of the plate thickness to 70 μm or less. The minimum grain size of the top 20% of grains with the largest grain size of bainite is preferably 60 μm or less, more preferably 55 μm or less, even more preferably 50 μm or less, and most preferably 45 μm or less. The lower limit is not particularly limited, but is preferably 10 μm or more, and more preferably 15 μm or more.

[0050] Here, the average grain size of bainite is determined as follows.

[0051] That is, after mirror polishing an L-section of a steel sheet (a section parallel to the rolling direction and parallel to the normal direction of the rolling surface), the crystal orientation of a randomly selected 1 mm x 1 mm region at the center of the sheet thickness was measured by electron backscatter diffraction (EBSD), and regions where the angle difference between adjacent pixels was 15° or more were determined as grain boundaries by image analysis.

[0052] The average grain size darea (μm) is calculated by dividing the area ai (μm 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 minimum grain size of the crystal grains when the circle-equivalent diameter of each crystal grain is used as the grain size and the largest 20% of the total number of crystal grains are selected.

[0053] At the center of the plate thickness, 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 were detected within an inspection area of ​​1 mm. ２ The reasons for limiting the size and 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.

[0054] Even if precipitates with a diameter of 100 nm or less (fine precipitates) are formed, the number of precipitates is 1 mm ２ If the number of precipitates per mm is less than 50,000, the effect of suppressing austenite grain growth is weak, and the decrease in strength in the intermediate 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.

[0055] 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).

[0056] (TS ０ -TS) / TS ０ In the present invention, the tensile strength (TS) at 350°C measured after aging under the condition of Larson Miller Parameter (LMP) = 15700 and the tensile strength (TS) at 350°C measured before aging are ０ ) is (TS ０ -TS) / TS０ ≦0.050. (TS ０ -TS) / TS ０ is an index for evaluating the decrease in tensile strength when held in the medium temperature range for a long time. If this index is 0.050 or less, the decrease in tensile strength after being held in the medium temperature range for a long time falls within a range that does not pose a problem in practical use. If this index is greater than 0.050, a significant decrease in tensile strength occurs after being held in the medium temperature range for a long time. For this reason, the above (TS ０ -TS) / TS ０ is 0.050 or less. (TS ０ -TS) / TS ０ is preferably 0.049 or less, 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.

[0057] 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). 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, most preferably 35 MPa or less, and most 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.

[0058] Aging treatment with an LMP of 15700 is performed under conditions of a heat treatment temperature and a heat treatment time such that the LMP, expressed by the following formula (2), becomes 15700. An LMP of 15700 corresponds to a condition of heat treatment at 350°C, which is a medium temperature range, for 20 years. The condition of Larson Miller Parameter (LMP) of 15700 refers to a condition 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) Furthermore, an example of an aging treatment condition under the above conditions is heat treatment at 400°C for 2335 hours.

[0059] Furthermore, 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. Although 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. Furthermore, the yield strength and tensile strength measured at 350°C after long-term aging in the intermediate 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 medium 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. Although there are no particular upper limits, the yield strength and tensile strength measured at 350°C after long-term aging in the medium 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 component composition to a specific value and employing the manufacturing conditions described below.

[0060] 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.

[0061] Toughness of weld heat affected zone: vE －４０ The toughness of the weld heat affected zone (HAZ) formed when the steel plate of the present invention is welded with the same steel plate or other steel is measured by a Charpy impact test at a test temperature of -40°C. －４０ It is 60J or more. －４０ If the absorbed energy vE is 60J or more, the toughness required for a structural pipe can be ensured. －４０ is preferably 70 J or more, more preferably 80 J or more, and even more preferably 100 J or more. There is no particular upper limit, but it is preferably 400 J or less, more preferably 380 J or less, and even more preferably 350 J or less. The notch position of the Charpy impact test specimen is 3 mm toward the base metal (HAZ 3 mm) from the bond portion, which is the boundary between the weld metal and the base metal. In addition, the absorbed energy (vE －４０) is considered to be within the scope of the present invention when the average value of the welding heat input is 60 J or more. 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.

[0062] <Steel Pipe> 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 high-strength welded steel pipe for steam transportation, even when it has a large diameter. The term "large diameter" means that the outer diameter (diameter) of the steel pipe (also referred to as "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. The maximum outer diameter is not particularly limited, but may be 1500 mm or less, and is 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 pipe for steam transportation.

[0063] The thickness of the steel pipe is not particularly limited, but is 12 to 30 mm when used for transporting steam. 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.

[0064] <Method of Manufacturing Steel Plate> Next, a method of manufacturing a steel plate according to the present invention will be described. The method of 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. 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 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.

[0065] Heating Process In the steel plate manufacturing method of the present invention, the heating process is a process 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 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. On the other hand, if the heating temperature exceeds 1200°C, austenite grain growth is significant, resulting in a deterioration in the toughness of the base material. Therefore, 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.

[0066] Hot Rolling Step In the method for producing a steel sheet of the present invention, the hot rolling step is a step of hot rolling the steel material heated in the heating step, including one or more passes of rolling at 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 at a rolling finish temperature of 850°C or less. The above hot rolling step is an important manufacturing condition of the present invention. By performing predetermined rolling in a temperature range of 900°C or less and setting the rolling finish temperature to 850°C or less, austenite grains are elongated, becoming finer in the thickness and width directions, and the dislocation density 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. As a result, strength, particularly in the medium temperature range, increases, and toughness is significantly improved.

[0067] If the cumulative reduction rate at 900°C or below is less than 50% or the rolling finish temperature exceeds 850°C, the austenite grains are not sufficiently refined, and the increase in intragranular dislocations is small. As a result, strength and toughness in the intermediate temperature range deteriorate. Therefore, the cumulative reduction rate at 900°C or below is set to 50% or more. The cumulative reduction rate is preferably 55% or more, more preferably 60% or more, even more preferably 65% ​​or more, and most preferably 70% or more.

[0068] The rolling end temperature is set to 850° C. or lower. In order to increase the reduction in the completely unrecrystallized region and achieve a finer structure, 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.

[0069] Although there is no particular upper limit to the cumulative reduction rate, an excessively large cumulative reduction rate may impose an excessive load on the rolling mill, so the cumulative reduction rate is preferably 95% or less, more preferably 90% or less, and even more preferably 85% or less. Furthermore, in order to make the bainite fraction 80% or more, the rolling end temperature is set to Ar ３ Temperatures above this are preferred. 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. 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 Just find the temperature.

[0070] Furthermore, by including one or more passes of rolling at 950°C or higher with a reduction rate per pass of 10% or more, austenite grains are refined by recrystallization, thereby improving the low-temperature toughness of the steel sheet. Therefore, it is necessary to include one or more passes of rolling at 950°C or higher with a reduction rate per pass of 10% or more. Furthermore, it is preferable to include two or more passes of rolling at 950°C or higher with a reduction rate per pass of 10% or more. It is more preferable to include three or more passes of rolling at 950°C or higher with a reduction rate per pass of 10% or more, and even more preferable to include four or more passes. Note that the upper limit of the number of passes of rolling at a reduction rate per pass of 10% or more 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. Furthermore, the upper limit of the reduction rate per pass at 950°C or higher is not particularly limited, but is preferably 40% or less.

[0071] Furthermore, by including one or more passes of rolling at 900°C or less with a reduction rate per pass of 15% or more, strain is introduced into the crystal grains, thereby promoting the refinement of austenite grains, thereby improving the low-temperature toughness of the steel sheet. Therefore, it is necessary to include one or more passes of rolling at 900°C or less with a reduction rate per pass of 15% or more. Furthermore, it is preferable to include two or more passes of rolling at 900°C or less with a reduction rate per pass of 15% or more. It is more preferable to include three or more passes of rolling at 900°C or less with a reduction rate per pass of 15% or more, and even more preferable to include four or more passes. Note that the upper limit of the number of passes of rolling at a reduction rate per pass of 15% or more 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. Furthermore, it is preferable to include one or more passes of rolling at 900°C or less with a reduction rate per pass of 20% or more. The upper limit of the rolling reduction per pass at 900° C. or less is not particularly limited, but it is preferably 40% or less.

[0072] Accelerated Cooling Step In the steel sheet manufacturing method of the present invention, 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. In this case, 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.

[0073] 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 needs to 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.

[0074] The strength of steel sheets tends to increase with increasing average cooling rate during accelerated cooling. When 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. Therefore, when the average cooling rate is less than 10°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 10°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 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. While there is no particular upper limit to the average cooling rate, to avoid an excessive increase in the martensite fraction, the average cooling rate is preferably 80°C / s or less, and more preferably 50°C / s or less.

[0075] 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, the growth of carbides 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. On the other hand, if the cooling stop temperature is lower than 250°C, the precipitation of low-temperature transformation products becomes significant, deteriorating the toughness of the base material, and the decomposition of the low-temperature transformation products in the mid-temperature range significantly reduces the strength in the mid-temperature range. Therefore, 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.

[0076] Reheating Process 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.

[0077] The reheating step after the accelerated cooling step, 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 the present invention. This process allows the precipitation of fine precipitates during reheating, which contributes to strengthening at room temperature and in the intermediate temperature range. To obtain fine precipitates, the material must be reheated to a temperature range of 550-700°C immediately after accelerated cooling. It is not necessary to set a specific temperature holding time during the reheating step. Furthermore, because precipitation progresses along with the bainite transformation during the cooling process after reheating, air cooling (including air blast cooling) is used after reheating. If the heating rate during the reheating step after accelerated cooling is less than 1°C / s, it takes a long time to reach the target reheating temperature, which reduces manufacturing efficiency. Furthermore, if the heating rate is less than 1°C / s, the precipitates grow, preventing the dispersion of fine precipitates and resulting in insufficient strength. Therefore, the temperature rise rate in the reheating step after accelerated cooling is set to 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. There is no particular upper limit to the temperature rise rate in the reheating step, but it is preferably 100° C. / s or less, more preferably 75° C. / s or less, and even more preferably 50° C. / s or less.

[0078] If the reheating temperature is less than 550°C, it falls outside the precipitation temperature range of Mo, Nb, and V, and sufficient precipitation strengthening cannot be achieved. Therefore, 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. On the other hand, if 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. Note that the heating rate of 1°C / s or higher in the reheating process 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, which 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.

[0079] Induction heating devices allow for easier temperature control than soaking furnaces and are relatively inexpensive. Furthermore, induction heating devices are particularly preferred because they can quickly heat steel sheets after cooling. 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 the type and size of steel sheets differ, simply by setting the number of induction heating devices to be energized and the power supply.

[0080] <Method for Producing Steel Pipe> The method for producing a steel pipe of the present invention includes a cold forming step, a welding step, and a pipe expansion step.

[0081] Cold-forming process The cold-forming process is a process of cold-forming the 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. It is also preferably 30 mm or less. The more preferred range is as described above. The method for cold-forming the steel plate into a tubular shape is not particularly limited. Examples of forming methods include cold-forming methods such as the UOE process, press bending (also called bending press), and roll forming to form a steel pipe. In the UOE process, grooves are formed on the widthwise ends of the steel plate, and then the widthwise ends of the steel plate are bent using a press. Subsequently, the steel plate is 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. In the case of press bending, the steel plate is successively formed by repeatedly bending the steel plate three-point. This produces a steel pipe having a substantially circular cross-sectional shape.

[0082] Welding Process 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. In this 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.

[0083] Pipe expansion process Next, after seam welding, the pipe is expanded to remove welding residual stress and improve the roundness of the steel pipe. In the pipe expansion process, 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.

[0084] The heat treatment after the production of the steel pipe may be carried out depending on the desired properties, and is not particularly specified.

[0085] Steel plates (steel symbols A to AM) having the chemical compositions shown in Table 1 were prepared under the manufacturing conditions shown in Tables 2-1 and 2-2 to have the thicknesses shown in Tables 2-1 and 2-2. These plates were then cold-formed, followed by submerged arc welding with a welding heat input of approximately 35 to 60 kJ / cm. The plates were then expanded at a 1.0% expansion ratio to produce steel pipes with an outer diameter of 610 mm and a pipe thickness of 15 to 25 mm. Note that, in the manufacturing conditions shown in Tables 2-1 and 2-2, "reduction ratio" refers to the cumulative reduction ratio at 900°C or less, "FT" refers to the rolling finish temperature, and "heat treatment" refers to long-term aging treatment. A 1.0% expansion ratio means that the inner diameter of the steel pipe is expanded by 1.0% 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.

[0086] A sample for steel structure observation 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 cross section parallel to the rolling direction and parallel to the normal direction of the rolling surface) was mirror-polished.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).The region where the angle difference between adjacent pixels was 15° or more was considered to be a grain boundary, and the crystal 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 μm.

[0087] The average grain size darea (μm) is calculated by dividing the area ai (μm 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 sizes 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 are selected.

[0088] In addition, 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 and then subjected to nital etching to reveal the microstructure.－２ mm ２ Steel structure photographs were taken of five fields of view (magnification: 400x), and 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 taken as average values ​​for 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.

[0089] The number of precipitates was determined using the evaluation method used in the embodiment, as described above.

[0090] To evaluate the low temperature toughness of the steel plate, a DWTT test piece was taken so that the longitudinal direction of the DWTT test piece was horizontal to the rolling direction of the steel plate and perpendicular to the plate thickness direction, and a DWTT test was carried out in accordance with API 5L to evaluate the ductile fracture surface area ratio and fracture surface transition temperature at a test temperature of -40°C. A case where the ductile fracture surface area ratio was 85% or more and the fracture surface transition temperature was -40°C or less at a test temperature of -40°C was evaluated as good.

[0091] For the steel sheet properties, 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 performed at 350°C to determine the yield strength and tensile strength. In the tensile test, round bar test specimens with a diameter of 6 mm were used, and a crosshead speed of 0.15 mm / min was used. A yield strength of 555 MPa or more and a tensile strength of 620 MPa or more at 350°C were evaluated as good. The steel sheet properties were evaluated by taking test specimens from the steel sheet before being formed into a steel pipe. For the steel pipe properties, round bar test specimens with a diameter of 6 mm were taken so that the longitudinal direction was the circumferential direction of the pipe, and a tensile test was performed at 350°C at a crosshead speed of 0.15 mm / min. The yield strength and tensile strength were evaluated in the same manner as for the steel sheet.

[0092] In addition, to simulate the high-temperature strength after long-term holding in the intermediate temperature range, the yield strength and tensile strength at 350°C were determined after heat treatment (aging treatment) in an N (nitrogen) atmosphere furnace under conditions of 15700 Larson-Miller Parameter (400°C, 2335 hours), which is the tempering parameter shown in equation (2), corresponding to 20 years of holding 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 the steel pipe in the same manner as before heat treatment. The results are shown in Tables 3-1 and 3-2. In addition, in order to evaluate the decrease in tensile strength when held in the intermediate temperature range for a long time, the tensile strength of the steel pipe properties (tensile strength before heat treatment (TS ０ )) - (tensile strength after heat treatment (TS))) / tensile strength before heat treatment (TS ０ ) was calculated, and a value of 0.050 or less was evaluated as being good.

[0093] In addition, in order to evaluate the decrease in yield strength when held in the medium temperature range for a long time, the difference in yield strength before and after the long-term aging was evaluated.

[0094] The toughness of the weld heat affected zone (HAZ) was evaluated by a Charpy impact test. The notch position of the Charpy impact test specimen was set at a position 3 mm (HAZ 3 mm) from the bond portion, which is the boundary between the weld metal and the base metal, toward the base metal side. The test was carried out at a temperature of -40°C. In the present invention, the Charpy impact test was carried out using three test specimens for each condition, and the absorbed energy (vE －４０ ) was evaluated as being excellent in toughness.

[0095] As mentioned above, Tables 2-1, 2-2, 3-1 and 3-2 show the manufacturing conditions of the steel plates and the test results of the steel plates and steel pipes.

[0096] Inventive Examples (Nos. 1 to 34) in which the chemical composition and the steel sheet manufacturing conditions are within the ranges of the present invention have a yield strength of 555 MPa or more and a tensile strength of 620 MPa or more before and after the heat treatment (measured at 350°C) of the steel sheet. In addition, in Inventive Examples (Nos. 1 to 34), the low temperature toughness of the steel sheet base material at -40°C, the HAZ toughness, and the (TS０ -TS) / TS ０ On the other hand, in the comparative examples (Nos. 35 to 67) in which the chemical composition or the steel sheet manufacturing conditions were outside the scope of the present invention, the yield strength, tensile strength, low temperature toughness at -40°C, HAZ toughness, and TS were all good before and after the heat treatment (measured at 350°C). ０ -TS) / TS ０ None of the targets were met.

[0097]

[0098]

[0099]

[0100]

[0101]

Claims

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 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 μm or less, and the minimum grain size of the top 20% of the bainite grains having the largest grain size is 70 μm or less, The yield strength before and after aging under the condition of Larson Miller Parameter (LMP) = 15700, as defined by the following formula (2), is 555 MPa or more, and the tensile strength (TS) at 350 ° C. measured after aging and the tensile strength (TS) at 350 ° C. measured before aging are different from each other. ０ ) is (TS ０ -TS) / TS ０ ≦0.050, the toughness of the steel plate is 85% or more in terms of the ductile fracture area ratio obtained by DWTT at -40 ° C, and the toughness of the weld heat affected zone formed during welding is vE －４０ Steel plate having a hardness of 60J or more. 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. LMP=(T+273)×(20+log(t))...(2) T: Heat treatment temperature (°C) t: heat treatment time (hours)

2. The steel sheet according to claim 1, wherein Ti / N is 2.0 to 4.0, and X represented by formula (3) is 0.70% or more. 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.

3. The component composition is expressed in mass% and further includes: Cu: 0.50% or less, Ni: 0.50% or less, Cr: 0.50% or less and Ca: Contains one or more of 0.0005 to 0.0040%, The steel sheet according to claim 1, wherein Y represented by formula (4) is 1.50% or less. Y=Cu+Ni+Cr+Mo...(4) The element symbols in formula (4) represent the content (mass%) of each element. Elements that are not contained are substituted with 0.

4. The component composition is in mass %, and further Cu: 0.50% or less, Ni: 0.50% or less, Cr: 0.50% or less and Ca: Contains one or more of 0.0005 to 0.0040%, The steel sheet according to claim 2, wherein Y represented by formula (4) is 1.50% or less. Y=Cu+Ni+Cr+Mo...(4) The element symbols in formula (4) represent the content (mass%) of each element, and 0 is substituted for elements that are not contained.

5. 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, At the center of the plate thickness, 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 were detected within an inspection area of ​​1 mm. ２ 5. The steel plate according to claim 1, wherein the number of particles is 50,000 to 1,000,000 per square inch, 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 MPa or less.

6. A steel pipe using the steel plate according to any one of claims 1 to 4.

7. A steel pipe using the steel plate described in claim 5.

8. The method for producing a steel sheet according to any one of claims 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 which, after cooling is stopped in the accelerated cooling step, the hot-rolled steel plate is reheated to obtain a steel plate under conditions of a heating rate of 1°C / s or higher and an ultimate temperature of 550 to 700°C within 150 seconds.

9. A method for manufacturing the steel plate according to claim 5, 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 which, after cooling is stopped in the accelerated cooling step, the hot-rolled steel plate is reheated to obtain a steel plate under conditions of a heating rate of 1°C / s or higher and an ultimate temperature of 550 to 700°C within 150 seconds.

10. A method for manufacturing a steel pipe, comprising: a cold forming step of cold-forming the steel plate according to any one of claims 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.

11. A method for manufacturing a steel pipe, comprising: a cold forming process for cold-forming the steel plate described in claim 5 into a tubular shape; a welding process for welding the butt joints where the ends of the steel plate formed into a tubular shape in the cold forming process are butted together; and an expansion process for expanding the pipe.