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

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

[Technical Field]

[0001] This invention relates to a steel sheet exhibiting excellent low-temperature toughness and a tensile strength of 620 MPa or higher after long-term aging in a medium temperature range, as well as a method for manufacturing the same, and a steel pipe made from the steel sheet and a method for manufacturing the same. In particular, the steel sheet of this invention is suitably usable as a material for high-strength steel pipes for steam piping. [Background technology]

[0002] There are two methods for recovering oil sands from oil reservoirs located in Canada and other countries: open-cut mining and steam injection, which involves introducing high-temperature, high-pressure steam into the reservoir through steel pipes. Because open-cut mining is only feasible in limited areas, steam injection is used in most regions. In steam injection, the steam introduced into the reservoir is in the 300-400°C temperature range (hereinafter referred to as the medium temperature range). In steam injection, this medium-temperature steam is introduced into the reservoir under high pressure. As mentioned above, steel pipes are used to introduce this steam. In recent years, with increasing energy demand, there has been a demand for larger diameter and higher strength steel pipes to improve the recovery rate of heavy oil and reduce installation costs. Furthermore, while installation and hydrostatic testing of these steel pipes are currently conducted in the warmer spring months, there is a growing demand for excellent low-temperature toughness to enable installation even in the colder winter months.

[0003] Patent documents 1 and 2 disclose a method for manufacturing steel pipes for steam transport 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 further increases in diameter are difficult for seamless steel pipes.

[0004] Furthermore, in the case of seamless steel pipes, a large amount of alloying elements are required to obtain a strength of API X80 grade or higher. In recent years, regarding manufacturing technology for high-strength steel pipes that can be manufactured by welding and have a large diameter, manufacturing technology for high-strength steel pipes having a strength of API X80 or higher has been disclosed in Patent Documents 3, 4, and 5. [Prior art documents] [Patent Documents]

[0005] [Patent Document 1] Japanese Patent Publication No. 2000-290728 [Patent Document 2] Patent No. 4821939 [Patent Document 3] Patent No. 5055736 [Patent Document 4] International Publication No. 2012 / 108027 [Patent Document 5] Patent No. 6137435 [Overview of the project] [Problems that the invention aims to solve]

[0006] Steel pipes manufactured using the method disclosed in Patent Document 3 satisfy the API X80 grade in high-temperature characteristics in the medium temperature range. However, Patent Document 3 does not disclose the strength characteristics of the steel pipes when used for a long period of time in the medium temperature range.

[0007] Patent Document 4 discloses a manufacturing technology for high-strength steel sheets with a yield strength of 700 MPa or higher. However, in order to ensure strength in the medium temperature range using the manufacturing technology for high-strength steel sheets disclosed in Patent Document 4, it is necessary to add a large amount of alloying components. Furthermore, it became clear during the process of developing the present invention that the high-strength steel sheets described in Patent Document 4 experience a significant decrease in tensile strength when held in the medium temperature range for a long period of time.

[0008] In the high-strength steel plate described in Patent Document 5, the high-temperature properties in the medium-temperature range satisfy the API X80 grade, and although the relationship (TS0 - TS) / TS0 ≤ 0.050 is satisfied for the tensile strength (TS) at 350°C measured after aging and the tensile strength (TS0) at 350°C measured before aging under the condition of LMP = 15700, the structure of the steel plate is not defined, and the low-temperature toughness is not considered either.

[0009] Thus, in the prior art, it is impossible to obtain a high-strength steel pipe for steam piping that satisfies all of the following: being large-diameter, having the strength characteristics required for high-strength steel pipes for steam transportation, and having excellent low-temperature toughness.

[0010] Steel pipes of API X80 or higher are required to satisfy the specifications of a yield strength of 555 MPa or more and a tensile strength of 620 MPa or more in a tensile test. Further, as characteristics required for high-strength steel pipes for steam transportation other than the above, in addition to satisfying the above strength characteristics before and after long-term aging in the medium-temperature range, being large-diameter, having excellent low-temperature toughness, and having excellent toughness in the heat-affected zone formed during welding can be mentioned.

[0011] The present invention has been made to solve the above problems, and the steel plate used as the material for the above steel pipe needs to satisfy all of the following: satisfying the specifications of a yield strength of 555 MPa or more and a tensile strength of 620 MPa or more in a tensile test at 350°C before and after aging, having excellent low-temperature toughness, and further having excellent toughness in the heat-affected zone formed during welding. An object of the present invention is to provide a steel plate used as the material for a steel pipe and a method for manufacturing the same. Another object of the present invention is also to provide a steel pipe composed of the above steel plate and a method for manufacturing the same.

[0012] Further, "having excellent low-temperature toughness" means that the ductile fracture surface rate: DWTTSA-40°C obtained by DWTT (test temperature: -40°C) conforming to API 5L is 85% or more, and the fracture transition temperature is -40°C or less. In DWTT, the test temperature was set to -40°C in anticipation of a decrease in toughness due to work hardening during pipe manufacturing.

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

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

Means for Solving the Problems

[0015] The inventors of the present invention have intensively studied the properties of steel sheets for large-diameter steel pipes in the medium-temperature range. As a result, by appropriately selecting the component composition and manufacturing conditions, it has been found that a steel sheet capable of manufacturing a steel pipe having the strength characteristics and low-temperature toughness required for high-strength steel pipes for steam transportation can be obtained while having a large diameter. In the manufacturing process of accelerated cooling after hot rolling and subsequent reheating, when reheating is performed during the bainite transformation of Nb-based steel in which Nb is dissolved, or Nb-V-based steel in which Nb and V are dissolved, in addition to the strengthening by bainite transformation during accelerated cooling, precipitation strengthening by fine precipitates precipitating from bainite and untransformed austenite during reheating, and suppression of dislocation recovery in the medium-temperature range make it possible to suppress the strength decrease in the medium-temperature range. In addition, when TiN is present, it becomes difficult for Nb to dissolve. As a result, compared with the case where Ti is not added, it becomes difficult for fine Nb carbides to disperse and precipitate during reheating after accelerated cooling, and it becomes difficult to suppress the strength decrease in the medium-temperature range. However, when the P eff value obtained by the following formula (1) is 0.050% or more, even in the case of adding Ti, sufficient dispersion precipitation of fine Nb carbides and V carbides during reheating can be obtained, and it becomes possible to suppress the strength decrease in the medium-temperature range. Nb carbides and V carbides refer to single carbides containing Nb or V, and composite carbides containing Nb and V. P eff (%) = (0.13Nb + 0.24V - 0.125Ti) / (C + 0.86N) ··· (1) In equation (1), the element symbols represent the mass percentage of each element present. For elements not present, 0 is substituted.

[0016] Furthermore, Nb and V are elements that form carbides in steel. Strengthening steel by precipitation of NbC has been a conventional practice. In addition, V-based carbides do not easily coagulate and coarse when held at high temperatures for a long time, making them useful elements for ensuring high-temperature creep strength. In this invention, the heating rate (temperature rise rate) during reheating after accelerated cooling is increased to suppress the growth of precipitates during heating. As a result, a large amount of finely precipitated Nb-containing carbides, or carbides containing both Nb and V, is formed in the steel, achieving an effect of suppressing strength reduction in the medium temperature range. In this invention, during reheating after accelerated cooling, the hot-rolled steel sheet is heated in an atmospheric furnace at a higher heating rate than that conventionally used in industry. By doing so, the growth of carbides containing Nb, or carbides containing both Nb and V, is suppressed, and a large amount of extremely fine precipitates with a particle size of less than 100 nm is obtained.

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

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

[0019] This invention was completed based on the above findings. Specifically, this invention provides the following: [1] In mass%, C: 0.040~0.090%, 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, The component composition contains N: 0.010% or less, with the remainder being Fe and unavoidable impurities, and the parameter P is represented by the following formula (1). eff The percentage is 0.050% or higher. The microstructure has bainite accounting for 80% or more of the area in the center of the plate 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 is 70 μm or less. The yield strength before and after aging, performed under the condition of Larson Miller Parameter (LMP) = 15700 as defined by equation (2) below, is 555 MPa or higher, the tensile strength (TS) measured at 350°C after aging and the tensile strength (TS0) measured at 350°C before aging satisfy the relationship (TS0-TS) / TS0 ≤ 0.050, the toughness of the steel plate is 85% or higher as determined by the ductile fracture surface ratio obtained by DWTT at -40°C, and the toughness of the heat-affected zone formed during welding is vE -40 Steel plates with a temperature of 60J or higher. P eff (%)=(0.13Nb+0.24V-0.125Ti) / (C+0.86N)...(1) In equation (1), the element symbols represent the mass percentage of each element present. For elements not present, 0 is substituted. LMP=(T+273)×(20+log(t))···(2) T: Heat treatment temperature (°C) t: Heat treatment time (hours) [2] The steel sheet described in [1], wherein the Ti / N ratio 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) In equation (3), the element symbols represent the mass percentage of each element. For elements that are not present, 0 is substituted. [3] The steel sheet according to [1] or [2], wherein the component composition is in mass%, and further contains one or more of the following: 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) In equation (4), the element symbols represent the mass percentage of each element. For elements that are not present, substitute 0. [4] The microstructure has ferrite with an area ratio of 10% or less and island-like martensite (MA) with an area ratio of 10% or less in 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 from Ti, Nb, V, Mo, Cr, and Al were observed in the center of the plate thickness over a 1 mm² area. 2 There are 50,000 to 1,000,000 of them, and the yield strength before aging determines the aging period. A steel plate as described in any of [1] to [3], wherein the difference in yield strength before and after aging, after subtracting the subsequent yield strength, is 50 MPa or less. [5] A steel pipe using the steel plate described in any of [1] to [4] above. [6] A method for manufacturing steel plates according to any of the above [1] to [4], A heating process in which the steel material is heated to 1000-1200°C, A hot rolling process in which the steel material heated in the heating step is hot-rolled under the conditions that the cumulative reduction ratio at 900°C or below is 50% or more, the reduction ratio per pass at 950°C or above is 10% or more in one or more passes, and the reduction ratio per pass at 900°C or below is 15% or more in one or more passes, and the rolling completion temperature is 850°C or below. The hot-rolled steel sheet obtained in the hot-rolling process is subjected to an accelerated cooling process 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. A method for manufacturing a steel sheet, comprising: a reheating step in which, after cooling is stopped in the accelerated cooling step, the hot-rolled steel sheet is reheated within 150 seconds under conditions of a heating rate of 1°C / s or more and a target temperature of 550 to 700°C to obtain a steel sheet. [7] A method for manufacturing a steel pipe, comprising: a cold forming step of cold forming a steel plate described in any of [1] to [4] into a tubular shape; a welding step of welding a butt joint formed by butting the ends of the steel plates formed into a tubular shape in the cold forming step; and a pipe expansion step of expanding the pipe. [Effects of the Invention]

[0020] According to the present invention, it is possible to obtain a steel pipe that, despite its large diameter, has a yield strength of 555 MPa or higher, a tensile strength of 620 MPa or higher, and excellent low-temperature toughness, all of which are required for high-strength steel pipes used for steam transport, after being held for a long period of time in a medium temperature range. Furthermore, according to the present invention, it is possible to obtain a steel pipe with the above characteristics even while reducing the amount of alloying elements used and thus reducing manufacturing costs. [Modes for carrying out the invention]

[0021] Embodiments of the present invention will be described below. However, the present invention is not limited to the embodiments described below.

[0022] <Steel plate> The steel sheet of the present invention (also called high-strength steel sheet) contains, by mass%, C: 0.040-0.090%, 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, and N: 0.010% or less. In the following description, "%" representing the content of components means "mass%".

[0023] C: 0.040~0.090% Carbon (C) is an essential element for ensuring the strength of steel through solid solution strengthening and precipitation strengthening. In particular, increasing the amount of solid solution C and forming precipitates contribute to ensuring strength in the medium temperature range. In order to ensure a predetermined strength at room temperature and in the medium temperature range, the present invention sets the C content to 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, if the C content exceeds 0.090%, it leads to deterioration of 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: 0.03~0.30% 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 should be 0.03% or more. Preferably, the Si content should be 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 deteriorates. Therefore, the Si content should be 0.30% or less. Also, preferably the Si content is 0.28% or less. More preferably, the Si content is 0.25% or less, even more preferably 0.23% or less, even more preferably 0.20% or less, and most preferably 0.18% or less.

[0025] Mn: 1.50~2.50% Mn is an effective element for improving the strength and toughness of steel. This effect can be fully obtained by increasing the Mn content to 1.50% or more. A Mn content of 1.55% or more is preferable. A Mn content of 1.60% or more is more preferable, 1.65% or more is even more preferable, 1.70% or more is even more preferable, and 1.75% or more is most preferable. Furthermore, if the Mn content exceeds 2.50%, toughness and weldability deteriorate significantly. Therefore, the Mn content should be 2.50% or less. Also, a Mn content of 2.40% or less is preferable, and 2.30% or less is more preferable. A Mn content of 2.20% or less is even more preferable, and 2.10% or less is most preferable.

[0026] P:0.020% or less P is an impurity element that significantly degrades toughness. Therefore, it is desirable to reduce the P content as much as possible. However, excessive reduction of the P content leads to increased manufacturing costs. Therefore, as a condition in which toughness degradation remains within an acceptable range, the P content should be 0.020% or less. Preferably, the P content should be 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 it is more preferably 0.003% or more.

[0027] S: 0.002% or less S is an impurity element and can significantly degrade toughness. Therefore, it is desirable to reduce the S content as much as possible. Furthermore, even if the morphology of S is controlled by adding Ca to change MnS to CaS-based inclusions, in the case of high-strength steel sheets of X80 grade or higher, finely dispersed CaS-based inclusions can also be a factor in toughness degradation. Therefore, the S content should be 0.002% or less. Preferably, the S content should be 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 it is preferably 0.0003% or more.

[0028] Mo: 0.10~0.60% Mo contributes significantly to increasing strength at room temperature and in the intermediate temperature range through solid solution or precipitate formation. However, if the Mo content is less than 0.10%, sufficient strength cannot be obtained in the intermediate temperature range, so it should be included at 0.10% or more. 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, if the Mo content exceeds 0.60%, toughness and weldability deteriorate, so it should be kept at 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: 0.020~0.070% Nb is an important element in this invention. Specifically, Nb is a necessary component for forming carbides and ensuring strength at room temperature and in the medium temperature range. Furthermore, Nb is necessary to refine the microstructure and impart sufficient strength and toughness by suppressing grain growth during slab heating and rolling. This effect is significant when the Nb content is 0.020% or more, so the Nb content should be 0.020% or more. Preferably, the Nb content should be 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 almost saturate, but the toughness deteriorates, so the Nb content should be 0.070% or less. Preferably, the Nb content should be 0.068% or less. More preferably, the Nb content should be 0.065% or less, even more 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 the heat-affected zone during welding. Ti has the effect of improving toughness by refining the microstructure. For this reason, 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 for fine carbides to disperse and precipitate, making it difficult to suppress the decrease in strength in the medium 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 composite precipitates with Ti and Nb, contributing to increased strength. Furthermore, V-based carbides are less prone to aggregation and coarsening even when held at high temperatures for extended periods, making V a useful element for ensuring high-temperature creep strength. If the desired high-temperature creep strength can be obtained by including elements other than V, the steel sheet of the present invention does not need to contain V, and the lower limit of the V content is not particularly limited, and may be 0%. To obtain 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 heat-affected zone deteriorates. Therefore, V content... The amount of nitrate is stipulated to be 0.080% or less. Furthermore, the V content should preferably be 0.070% or less. It is preferable that the V content be 0.060% or less, even more preferably 0.050% or less, and most preferably 0.040% or less.

[0032] Al: 0.045% or less Al is added as a deoxidizing agent. To obtain the deoxidizing effect, 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 its toughness deteriorates. Therefore, the Al content should be 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 reaching 1350 °C or higher. Due to this fine dispersion, the prior austenite grains in the weld heat-affected zone are refined and the toughness of the weld heat-affected zone is improved. To obtain this effect, it is preferable that the N content is 0.002% or more, and more preferably 0.0025% or more. Also, when the N content exceeds 0.010%, the toughness of the base metal deteriorates due to coarsening of the precipitates and increase in dissolved N, and furthermore, the toughness of the weld metal when made into a steel pipe deteriorates. Therefore, the N content is 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 by (0.13Nb + 0.24V - 0.125Ti) / (C + 0.86N) ··· (1). In this formula (1), the element symbols mean the content (% by mass) of each element, and 0 is substituted for elements not contained. P eff It is necessary in the present invention to adjust the content of the above elements so that P becomes 0.050% or more. eff P is an important factor for making the steel composed of the above composition range have excellent strength in the medium temperature range. eff When (%) is less than 0.050%, the amount of finely dispersed carbides precipitated during reheating after cooling decreases. As a result, the strength, particularly the tensile strength after long-time heat treatment, significantly decreases. Therefore, P eff (%) is 0.050% or more. eff P is preferably 0.055% or more in order to sufficiently suppress the strength decrease after heat treatment. eff (%) is more preferably 0.060% or more, even more preferably 0.065% or more, and most preferably 0.070% or more. Also, in order to suppress the generation of a large amount of precipitates in the weld heat-affected zone and deterioration of toughness, P effIt 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] The steel sheet of the present invention may further contain one or more of Cu, Ni, Cr, and Ca in addition to the above-mentioned component composition, for the purpose of further improving its properties.

[0036] Cu: 0.50% or less Cu is one of the elements effective in improving toughness and increasing strength. To obtain this effect, it is preferable to have a Cu content of 0.05% or more, and more preferably 0.10% or more. On the other hand, a Cu content exceeding 0.50% inhibits weldability, so if Cu is included, the Cu content should be 0.50% or less. It is preferable to have a Cu content of 0.48% or less. It is more preferable to have a Cu content of 0.45% or less, even more preferable to have a Cu content of 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, if Ni is included, its content should be 0.50% or less. The Ni content is preferably 0.48% or less. The Ni content is more preferably 0.45% or less, even more preferably 0.43% or less, and most preferably 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. If the Cr content exceeds 0.50%, it adversely affects weldability. Therefore, when Cr is included, the Cr content should be 0.50% or less. The Cr content is preferably 0.48% or less, more preferably 0.45% or less, even more preferably 0.43% or less, and most preferably 0.40% or less.

[0039] Ca: 0.0005~0.0040% Ca controls the morphology of sulfide inclusions and improves toughness. This effect becomes apparent when the Ca content is 0.0005% or higher. Therefore, when Ca is included, the Ca content should be 0.0005% or higher. Preferably, the Ca content should be 0.0010% or higher. If the Ca content exceeds 0.0040%, not only does the above effect saturate, but the cleanliness decreases and the toughness deteriorates. Therefore, when Ca is included, its content should be 0.0040% or less. Preferably, the Ca content is 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: 2.0~4.0 (preferred requirement) By specifying the Ti / N ratio within an appropriate range, TiN is finely dispersed, achieving refinement of the prior austenite grains in the heat-affected zone (HZ). This refinement of the prior austenite grains in the HZ improves the toughness of the HZ in the low-temperature range below -40°C and the medium-temperature range above 300°C. If the Ti / N ratio is less than 2.0, the effect is not sufficient, so it is preferable to set the Ti / N ratio to 2.0 or higher, more preferably 2.1 or higher, even more preferably 2.2 or higher, and most preferably 2.3 or higher. If the Ti / N ratio exceeds 4.0, it leads to coarsening of the prior austenite grains due to the coarsening of precipitates. Since the toughness of the HZ deteriorates due to the coarsening of the prior austenite grains, it is preferable to set the Ti / N ratio to 4.0 or lower, more preferably 3.9 or lower, even more preferably 3.8 or lower, and most preferably 3.7 or lower.

[0041] X: 0.70% or more (preferred requirement) X=0.35Cr+0.9Mo+12Nb+8V...(3) In equation (3), the element symbols represent the mass percentage of each element. For elements that are not present, 0 is substituted. Cr, Mo, Nb, and V contribute to improving tempering softening resistance and strengthening of intragranular precipitation during rolling in steel composed within the above component range. Equation (3) above is an important factor in obtaining steel that has excellent strength of X80 grade or higher in the medium temperature range after long-time heat treatment, and also has good low-temperature toughness. The effect of satisfying equation (3) is greatly realized when combined with the manufacturing conditions described later. In this invention, it is preferable that X be 0.70% or more in order to achieve X80 grade strength after long-time heat treatment at 350°C. More preferably, X be 0.75% or more. In order to achieve X100 grade strength after long-time heat treatment at 350°C, it is even more preferable that X be 0.90% or more. Most preferably, it be 1.00% or more. Furthermore, if X is 2.00% or more, the low-temperature toughness of the weld may decrease. Therefore, it is preferable that X be 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) In equation (4), the element symbols represent the amount of each element present, and 0 is substituted for elements that are not present. In formula (4) above, Y is preferably 1.50% or less. To balance the effect of increasing strength due to these elements with manufacturing costs, it is preferable to have a total content of the above elements of 1.50% or less. 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 it is preferable that Y be 0.20% or more, and more preferably 0.30% or more. It should be noted that one of the features of the present invention is that desired properties can be obtained even when the amount of these components used is reduced.

[0043] The remainder of the components other than those listed above consists of Fe and unavoidable impurities. These impurities are inevitably introduced from raw materials, manufacturing processes, or manufacturing equipment, and their inclusion is permissible as long as it does not hinder the objectives of the present invention. Examples of raw materials include iron ore, reduced iron, or scrap. Examples of unavoidable 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: Over 80% by area Bainite is an important microstructure for achieving both strength and low-temperature toughness. Furthermore, bainite effectively contributes to improving the strength of steel sheets through transformation strengthening. In addition, uniformity of the microstructure is necessary from the viewpoint of increasing the initial dislocation density and improving the strength of high-strength steel sheets, especially in the medium temperature range. For this reason, the microstructure of the steel sheet of the present invention must be predominantly bainite, and specifically, bainite must account for 80% or more of the total area of ​​the steel microstructure at the center of the sheet thickness. For the reason of ensuring a strength-toughness balance, bainite is preferably at 83% or more, more preferably at 85% or more, even more preferably at 88% or more, and most preferably at 90% or more. There is no particular upper limit to the bainite fraction, but from the viewpoint of improving deformation performance, bainite is preferably at 98% or less, and more preferably at 95% or less.

[0046] Ferrite: Area ratio of 10% or less (preferred conditions) When ferrite is processed by rolling, the number of movable dislocations increases, which can raise its strength. On the other hand, heat treatment at medium temperatures significantly reduces the dislocation density of the ferrite phase, causing it to soften and potentially reducing the strength of the steel sheet after heat treatment. Therefore, it is preferable that the area ratio of ferrite at the center of the sheet thickness be 10% or less. More preferably, the area ratio of ferrite is 9% or less, even more preferably 8% or less, and most preferably 7% or less. There is no particular lower limit to the area ratio of ferrite, but from the viewpoint of improving deformation performance, it is preferable that the area ratio of ferrite be 1% or more, and more preferably 2% or more.

[0047] Island-like martensite (MA): Area ratio of 10% or less (preferred conditions) Island martensite (MA: Martensite-Austenite constituent) is a very hard phase and can act as a fracture initiation point, potentially reducing the low-temperature toughness of steel sheets. Therefore, it is preferable that the area ratio of island martensite (MA) in the center of the sheet thickness be 10% or less. More preferably, the area ratio of island martensite is 8% or less, even more preferably 5% or less, and most preferably 3% or less. There is no particular lower limit to the island martensite fraction, but it is preferably 1% or more in area ratio, and preferably 2% or more in area ratio. Furthermore, as mentioned above, the microstructure of the base steel sheet must basically consist of the bainite described above, but other microstructures besides bainite, ferrite, and island martensite (MA) include pearlite, martensite, cementite, and retained austenite.

[0048] 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 is 70 μm or less. Since the grain boundaries of bainite provide resistance to brittle crack propagation, grain refinement contributes to improved low-temperature toughness. Therefore, the average grain size of bainite should be 30 μm or less. Preferably, the average grain size of bainite should be 28 μm or less. More preferably, the average grain size of bainite should be 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 5 μm or more is preferred, and 6 μm or more is more preferred.

[0049] Furthermore, while refining the average grain size improves low-temperature toughness, there are limitations to refining the average grain size when considering plate crowns or when cooling is initiated from high temperatures. Moreover, in this invention, it is crucial to suppress the formation of coarse grains. Coarse bainite is highly likely to be the starting point for fracture, and if the minimum grain size of the top 20% of grains with large grain size is large, low-temperature toughness deteriorates. In particular, regarding the grain size of bainite, if the minimum grain size of the top 20% of grains with large grain size exceeds 70 μm, it is likely to become the starting point for fracture. Therefore, it is necessary to keep the minimum grain size of the top 20% of grains with large bainite grain size in the center of the plate thickness 70 μm or less. The minimum grain size of the top 20% of grains with large bainite grain size 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. There is no particular limit to the lower limit, but it is preferably 10 μm or more, and more preferably 15 μm or more.

[0050] Here, the average grain size of bainite was determined as follows:

[0051] Specifically, the L-shaped cross-section of the steel plate (a cross-section parallel to the rolling direction and parallel to the direction normal to the rolling surface) was mirror-polished, and the crystal orientation of a randomly selected 1 mm × 1 mm region at the center of the plate thickness was measured using electron beam backscatter diffraction (EBSD). The region where the angle difference between adjacent pixels is 15° or more was identified as a grain boundary and determined by image analysis.

[0052] Note that the average crystal grain size darea (μm) is the area ai (μm) occupied by each crystal grain. 2 ) and the equivalent circular diameter di (μm) of each crystal grain were calculated using the following formula. darea(μm)=Σ(ai·di) / Σai Furthermore, the minimum grain size for the top 20% of bainite grains is calculated by arranging the grains in descending order of their circular diameter (equivalent to the grain size) and selecting 20% ​​of the total grains from largest to smallest.

[0053] 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 from Ti, Nb, V, Mo, Cr, and Al were observed in the center of the plate thickness over a 1 mm² area. 2 Approximately 50,000 to 1,000,000 exist per unit area (optimal conditions). This document explains 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 from Ti, Nb, V, Mo, Cr, and Al (for example, one or more of Ti-Nb precipitates, Ti-V precipitates, Cr-Mo precipitates, Cr-Mo-Nb precipitates, Al-Ti precipitates, Al-Nb precipitates, and Al-Ti-Nb precipitates) in the steel of the present invention.

[0054] Even if precipitates with a diameter of 100 nm or less (micro-precipitates) are formed, if the number of precipitates is 1 mm 2 If the number of precipitates per unit area is less than 50,000, the austenite grain growth inhibitory effect is weak, and the decrease in strength in the medium temperature range may not be suppressed. Therefore, the number of precipitates with a diameter of 100 nm or less in the steel at the center of the plate thickness should be 1 mm. 2 It is preferable to have 50,000 or more particles per unit area. More preferably 80,000 or more, even more preferably 100,000 or more, and most preferably 130,000 or more. On the other hand, precipitates with a diameter of 100 nm or less are 1 mm2 If the number exceeds 1,000,000 per unit area, the aggregation and coarsening of the fine precipitates can conversely weaken the austenite grain growth inhibitory effect, potentially leading to a greater decrease in strength in the medium temperature range. Therefore, precipitates with a diameter of 100 nm or less should be 1 mm 2 It is preferable to have 1,000,000 or fewer per unit. More preferably 950,000 or fewer, even more preferably 900,000 or fewer, and most preferably 850,000 or fewer.

[0055] Here, we will explain the method for measuring the density and size (diameter) of precipitates. Observe the corrosion surface at any location on the head cross-section using a scanning electron microscope (SEM), or prepare an extraction replica sample or thin film sample and observe it using a transmission electron microscope (TEM). The number of precipitates smaller than 100 nm is measured at least 100 μm. 2 Measurements are taken over the area specified above. For example, when observing with a magnification of 100,000x and a field of view of 2000nm x 2000nm, the observation area per field of view is 4μm². 2 Therefore, 25 fields of view are observed randomly. These measurement results are converted to the number of particles per unit area. The number of precipitates smaller than 100 nm is calculated for 25 fields of view (100 μm). 2 If there are 5 of them, the density of the precipitate is 1 mm 2 This can be converted to 50,000 particles per unit area. The density of the precipitate is measured over a 1 mm² area. 2 This is called the number of particles per unit. The diameter of the precipitate is defined as the average of the major axis (long side) and the minor axis (short side).

[0056] (TS0-TS) / TS0≦0.050 In this invention, 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 medium temperature range. If this index is 0.050 or less, the decrease in tensile strength after being held for a long time in the medium temperature range is within a range that does not pose a practical problem. If this index is greater than 0.050, a significant decrease in tensile strength occurs after being held for a long time in the medium temperature range. For this reason, (TS0-TS) / TS0 is set to 0.050 or less. (TS0-TS) / TS0 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 restricted and can be a negative value or greater than -0.100.

[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 this 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 time in the medium temperature range. If this difference is 50 MPa or less, the decrease in yield strength after being held for a long time in the medium temperature range will be within a range that does not pose a practical problem. If this difference is greater than 50 MPa, a significant decrease in yield strength will occur after being held for a long time in the medium 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, and may be -100 MPa or more.

[0058] Aging treatment with LMP = 15700 refers to aging treatment performed under heat treatment temperature and heat treatment time conditions that result in an LMP of 15700, as expressed by equation (2) below. LMP = 15700 corresponds to heat treatment at a medium temperature of 350°C for 20 years. Note that the condition for Larson Miller Parameter (LMP) = 15700 means that it is between 15650 and 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 higher and a tensile strength of 620 MPa or higher, measured at 350°C. The yield strength and tensile strength measured at 350°C were both 560 MPa. It is preferable that the yield strength and tensile strength measured at 350°C be 840 MPa or less and 900 MPa or less, respectively, and more preferably 830 MPa or less and 890 MPa or less, respectively. Furthermore, it is preferable that the yield strength and tensile strength measured at 350°C be 840 MPa or less and 900 MPa or less, respectively, and more preferably 830 MPa or less and 890 MPa or less, respectively. In addition, the yield strength measured at 350°C after long-term aging in the medium temperature range is 555 MPa or more and the tensile strength is 620 MPa or more. It is preferable that the yield strength and tensile strength measured at 350°C after long-term aging in the medium temperature range be 560 MPa or more and 625 MPa or more, respectively, more preferably 565 MPa or more and 630 MPa or more, respectively, and 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 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 specific component composition and adopting the manufacturing conditions described later.

[0060] Toughness of steel plate: Ductile fracture surface ratio obtained by DWTT at -40°C is 85% or higher. The toughness of the steel sheet of the present invention is such that the ductile fracture surface ratio (DWTTSA-40℃) obtained by DWTT (test temperature: -40℃) in accordance with API 5L is 85% or higher. If the above value is less than 85%, the steel sheet will undergo brittle fracture at low temperatures, making it difficult to lay the sheet throughout the year, including winter when temperatures fall below 0℃, or to use it in areas with very low ambient temperatures. Therefore, it is necessary to obtain the above value. It also means that the fracture transition temperature is -40℃ or lower. The reason for setting the test temperature to -40℃ in DWTT is to account for the decrease in toughness due to work hardening during pipe manufacturing. Furthermore, the ductile fracture surface ratio is preferably 86% or higher, more preferably 87% or higher, and even more preferably 88% or higher. The upper limit is not particularly limited, and the ductile fracture surface ratio obtained by DWTT at -40℃ may be 100% or less.

[0061] Toughness of the area affected by welding heat: vE -40 60J or more The toughness of the heat-affected zone (HAZ) formed when the steel sheet of the present invention is welded to the same steel sheet or another steel sheet is determined by the absorbed energy vE when measured by a Charpy impact test at a test temperature of -40°C. -40 It is 60J or more. vE -40 If the absorbed energy vE is 60J or higher, the toughness required for a structural pipe can be secured. -40 It 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, and 3 A value of 80J or less is more preferable, and 350J or less is even more preferable. The notch position of the Charpy impact test specimen is 3 mm (HAZ 3 mm) from the bond area, which is the boundary between the weld metal and the base metal, towards the base metal. Furthermore, the absorbed energy (vE) when the Charpy impact test was performed using three test specimens for each condition is also specified. -40The present invention is defined as being within the range where the average value of ) is 60J or more. The upper limit is not particularly limited. The welding method shall be submerged arc welding, and the welding heat input shall be 85kJ / cm or less, depending on the thickness of the steel plate. Preferably, the heat input shall be 80kJ / cm or less, and more preferably 70kJ / cm or less. Furthermore, the lower limit is not particularly limited, but preferably 20kJ / cm or more, and more preferably 25kJ / cm or more.

[0062] <Steel pipe> Since the steel pipe of the present invention is manufactured using the steel plate of the present invention, even in large diameters, it possesses the strength characteristics and low-temperature toughness required for high-strength welded steel pipes used for steam transport. Large diameter means that the outer diameter (diameter) of the steel pipe (also called high-strength steel pipe) is 400 mm or more. Preferably, the outer diameter of the steel pipe is 500 mm or more, more preferably 600 mm or more, and even more preferably 700 mm or more. In particular, the maximum outer diameter is not limited, but may be 1500 mm or less, and more preferably 1400 mm or less. According to the present invention, the diameter can be increased while maintaining the strength characteristics required for high-strength welded steel pipes for steam transport.

[0063] Furthermore, the thickness of the steel pipe is not particularly limited, but in the case of steam transport, it is 12 to 30 mm. That is, the thickness of the steel pipe is preferably 12 mm or more, more preferably 13 mm or more, even more preferably 14 mm or more, and most preferably 15 mm or more. Also, 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 plates> Next, the method for manufacturing steel sheets according to the present invention will be described. The method for manufacturing steel sheets according to the present invention comprises a heating step, a hot rolling step, an accelerated cooling step, and a reheating step. Unless otherwise specified, the temperature in the description of each step shall be the average temperature in the thickness direction of the steel sheet. The average temperature in the thickness direction can be determined from the surface temperature of the slab (steel material) or steel sheet by simulation calculations using parameters such as sheet thickness, cooling conditions, and heat transfer coefficient. For example, the average temperature in the thickness direction can be determined by calculating the temperature distribution in the thickness direction using the 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 stop of cooling. The heating rate (temperature rise rate) in the reheating step is the average temperature rise rate obtained by dividing the temperature difference required to reheat to the reheating temperature after cooling in the accelerated cooling step by the time required to reheat.

[0065] heating process In the steel sheet manufacturing method of the present invention, the heating step is a step of heating the steel material to 1000 to 1200°C. Here, the steel material is, for example, a slab obtained by casting molten steel. Since the component composition of the steel material becomes the component composition of the steel sheet, the adjustment of the component composition of the steel sheet can be done at the stage of adjusting the component composition of the molten steel. There are no particular limitations on the steelmaking method of the steel material. In the hot rolling step described later, the heating temperature is set to 1000°C or higher in order to sufficiently advance austenitization and solid solution of carbides and to obtain sufficient strength at room temperature and in the medium 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, the growth of austenite grains is significant and the toughness of the base material deteriorates. Therefore, the heating temperature is set to 1200°C or lower. Furthermore, the heating temperature is preferably 1190°C or lower, more preferably 1180°C or lower, even more preferably 1170°C or lower, and most preferably 1160°C or lower.

[0066] Hot rolling process In the steel sheet manufacturing method of the present invention, the hot rolling step is a step in which the steel material heated in the heating step is hot-rolled under the conditions that the cumulative reduction rate at 900°C or below is 50% or more, the reduction rate per pass at 950°C or above is 10% or more in one pass, the reduction rate per pass at 900°C or below is 15% or more in one pass, and the rolling completion temperature is 850°C or below. The hot rolling process described above is an important manufacturing condition of the present invention. By performing a predetermined rolling process in a temperature range of 900°C or lower, and ending the rolling at a temperature of 850°C or lower, the austenite grains expand, becoming finer in the thickness and width directions of the plate, and the dislocation density within the grains introduced by rolling increases. This effect is achieved by including at least one rolling pass where the cumulative reduction ratio at temperatures below 900°C is 50% or more, and the reduction ratio per pass at temperatures above 950°C is 10% or more, and at least one rolling pass where the reduction ratio per pass at temperatures below 900°C is 15% or more, with the rolling completion temperature set to 850°C or below. As a result, strength, especially strength in the medium temperature range, increases and toughness is significantly improved.

[0067] If the cumulative reduction ratio at temperatures below 900°C is less than 50% or the rolling completion temperature exceeds 850°C, the austenite grains are not sufficiently refined, and the increase in dislocations within the grains is small. As a result, the strength and toughness in the intermediate temperature range deteriorate. Therefore, the cumulative reduction ratio at temperatures below 900°C should be 50% or more. A cumulative reduction ratio of 55% or more is preferable, 60% or more is more preferable, 65% or more is even more preferable, and 70% or more is most preferable.

[0068] Furthermore, the rolling completion temperature should be 850°C or lower. In order to increase the reduction amount in the completely unrecrystallized region and achieve microstructure refinement, the rolling completion temperature should preferably be 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] The above cumulative reduction ratio does not have any particular upper limit, but if the cumulative reduction ratio is made excessively large, there is a risk of placing an excessive load on the rolling mill, so it is preferable that the cumulative reduction ratio be 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 preferably the Ar3 temperature or higher. The Ar3 temperature can be calculated by the following formula. Ar3(℃)=910-310[C]-80[Mn]-20[Cu]-55[Ni]-15[Cr]-80[Mo] However, in the formula, [C], [Mn], [Cu], [Ni], [Cr], and [Mo] represent the content (mass%) of C, Mn, Cu, Ni, Cr, and Mo in the base steel sheet, respectively. If some elements are not present in the base steel sheet, the content of those elements should be set to "0" when calculating the Ar3 temperature.

[0070] Furthermore, including one or more rolling passes at 950°C or higher with a reduction ratio of 10% or more per pass promotes the recrystallization of austenite grains, thereby improving the low-temperature toughness of the steel sheet. For this reason, it is necessary to include one or more rolling passes at 950°C or higher with a reduction ratio of 10% or more per pass. It is also preferable to include two or more rolling passes at 950°C or higher with a reduction ratio of 10% or more per pass. It is more preferable to include three or more rolling passes at 950°C or higher with a reduction ratio of 10% or more per pass, and even more preferable to include four or more passes. There is no particular upper limit to the number of rolling passes at which the reduction ratio per pass is 10% or more, but it may be 50 passes or less, more preferably 30 passes or less, and even more preferably 15 passes or less. Furthermore, there is no particular upper limit to the reduction ratio per pass at 950°C or higher, but it is preferably 40% or less.

[0071] Furthermore, by including one or more rolling passes at temperatures below 900°C with a reduction ratio of 15% or more per pass, strain is introduced into the crystal grains, promoting the refining of austenite grains and thus improving the low-temperature toughness of the steel sheet. For this reason, it is necessary to include one or more rolling passes at temperatures below 900°C with a reduction ratio of 15% or more per pass. It is also preferable to include two or more rolling passes at temperatures below 900°C with a reduction ratio of 15% or more per pass. It is more preferable to include three or more rolling passes at temperatures below 900°C with a reduction ratio of 15% or more per pass, and even more preferable to include four or more passes. There is no particular upper limit to the number of rolling passes at which the reduction ratio per pass is 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. Furthermore, it is preferable to include one or more rolling passes at temperatures below 900°C with a reduction ratio of 20% or more per pass. Furthermore, while there is no particular upper limit to the reduction ratio per pass at temperatures below 900°C, it is preferable that it be 40% or less.

[0072] Accelerated cooling process In the steel sheet manufacturing method of the present invention, the accelerated cooling step is a step in which the hot-rolled steel sheet obtained in the hot-rolling step is cooled at a starting temperature of 700°C or higher, an average cooling rate of 10°C / s or higher, and a cooling stop temperature of 2°C / s. This process involves accelerated cooling under conditions of 50-550°C. In this case, the average cooling rate is defined as the cooling rate obtained by dividing the difference between the cooling start temperature and the cooling stop temperature by the time required from the start to the stop of cooling.

[0073] To suppress ferrite formation 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. Preferably, the cooling start temperature is 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. 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. Therefore, if the average cooling rate is less than 10°C / s, sufficient strength cannot be obtained at room temperature and in the medium temperature range. Also, if the average cooling rate is less than 10°C / s, the effect of refining the microstructure becomes smaller, the grain size does not decrease, and the low-temperature toughness deteriorates. For this reason, the average cooling rate during accelerated cooling should be 10°C / s or higher. The average cooling rate is preferably 12°C / s or higher, more preferably 14°C / s or higher, even more preferably 16°C / s or higher, and most preferably 18°C / s or higher. There is no particular upper limit to the average cooling rate, but in order to avoid an excessive increase in the martensite fraction, the average cooling rate is preferably 80°C / s or lower, and more preferably 50°C / s or lower.

[0075] The strength of steel plates 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 dissolved carbon decreases. As a result, sufficient strength, especially sufficient strength in the medium temperature range, cannot be obtained. For this reason, the cooling stop temperature should be 550°C or lower. Preferably, the cooling stop temperature should be 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 below 250°C, the precipitation of low-temperature transformation products becomes significant, the toughness of the base material deteriorates, and the strength in the medium temperature range decreases significantly due to the decomposition of low-temperature transformation products in the medium temperature range. For this reason, the cooling stop temperature of accelerated cooling should be 250°C or higher. Preferably, the cooling stop temperature should be 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 step is a process in which the hot-rolled steel sheet is reheated within 150 seconds (s) after the cooling stops in the accelerated cooling step, under conditions of a heating rate of 1°C / s or more and a target temperature of 550 to 700°C. The time from the cooling stop in the accelerated cooling step to the start of the reheating step shall be within 150 seconds, preferably within 130 seconds, more preferably within 120 seconds, even more preferably within 110 seconds, and most preferably within 100 seconds. There is no particular lower limit to the time from the cooling stop in the accelerated cooling step to the start of the reheating step, but it is preferably 5 seconds or more, more preferably 6 seconds or more, and even more preferably 7 seconds or more.

[0077] Accelerated cooling process performed with a heating rate of 1°C / s or more and a target temperature of 550-700°C. The subsequent reheating step is important in this invention. This process allows for the precipitation of fine precipitates during reheating, which contribute to strengthening at room temperature and in the medium temperature range. To obtain fine precipitates, it is necessary to reheat to a temperature range of 550-700°C immediately after accelerated cooling. There is no need to set a specific temperature holding time in the reheating step. Furthermore, since precipitation progresses along with bainite transformation during the cooling process after reheating, cooling after reheating should be done by air cooling (including blast air cooling). If the heating rate in the reheating step after accelerated cooling is less than 1°C / s, it will take a long time to reach the target reheating temperature, resulting in poor manufacturing efficiency. Also, if the heating rate is less than 1°C / s, the precipitates will grow, preventing the dispersion and precipitation of fine precipitates from being obtained, and thus insufficient strength cannot be obtained. Therefore, 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. Furthermore, while there is no particular upper limit to the heating rate in the reheating process, 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 below 550°C, the precipitation temperature range for Mo, Nb, and V will be outside this range, and sufficient precipitation strengthening will not be achieved. Therefore, the reheating temperature should be 550°C or higher. Preferably, the reheating temperature should be 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 precipitate will become coarse and sufficient strength cannot be obtained at room temperature and in the medium temperature range. Therefore, the reheating temperature should be 700°C or lower. Preferably, the reheating temperature should be 690°C or lower, more preferably 680°C or lower, even more preferably 670°C or lower, and most preferably 660°C or lower. Furthermore, achieving a heating rate of 1°C / s or more in the reheating process after accelerated cooling as defined in this invention is difficult in an atmospheric furnace depending on the plate thickness. Therefore, it is preferable to use a gas combustion furnace or induction heating device that can rapidly heat steel plates as the heating device. Moreover, it is even more preferable to install the gas combustion furnace or induction heating device on the conveyor line downstream of the cooling equipment used for accelerated cooling.

[0079] Induction heating devices offer easier temperature control and relatively lower costs compared to soaking furnaces and the like. Furthermore, induction heating devices are particularly preferable because they can quickly heat steel plates after cooling. By arranging multiple induction heating devices in series, the heating rate and reheating temperature can be freely controlled simply by setting the number of induction heating devices and the power supply, even when line speeds, steel plate types, or dimensions differ.

[0080] <Method of manufacturing steel pipes> The present invention provides a method for manufacturing steel pipes, comprising 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 sheet of the present invention into a tubular shape. When manufacturing steel pipes for steam transport, the thickness of the steel sheet is preferably 12 mm or more. It is also preferably 30 mm or less. The more preferable range is as described above. The method of forming the steel sheet into a tubular shape by cold forming 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 the steel pipe shape. In the UOE process, the widthwise ends of the steel sheet are beveled, then the widthwise ends of the steel sheet are bent using a press machine, and then the steel sheet is formed into a U-shape and then into an O-shape using a press machine, thereby forming the steel sheet into a cylindrical shape so that the widthwise ends of the steel sheet face each other. In the case of press bending, steel pipes with a nearly circular cross-sectional shape are manufactured by repeatedly bending a steel plate at three points in sequence.

[0082] Welding process The welding process is the process of welding the butt joints of steel plates that have been formed into a tubular shape in the cold forming process. The welding method is not particularly limited, but it is acceptable to weld them together by submerged arc welding or the like. The opposing ends in the width direction of the steel plates are butted together and welded. This type of welding is called seam welding. In this seam welding, a method having two stages is preferred: a tack welding process in which the cylindrical steel plates are restrained and the ends in the width direction of the steel plates are butted together and tack-welded, and a main welding process in which welding is performed on the inner and outer surfaces of the butt joint of the steel plates by submerged arc welding.

[0083] Pipe expansion process Next, after seam welding, the pipe is expanded to remove residual stress from welding and improve the roundness of the steel pipe. In the pipe expansion process, the expansion ratio (the change in the 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. Furthermore, it is preferable that the expansion ratio be 1.5% or less. Moreover, from the viewpoint of balancing the effect of improving roundness and the capacity required of the pipe expansion equipment, it is even more preferable that the expansion ratio be 0.5% or more. It is even more preferable that the expansion ratio be 1.2% or less.

[0084] Heat treatment after steel pipe manufacturing may be carried out according to the desired properties and is not specifically prescribed. [Examples]

[0085] Using steel materials (steel symbols A to AM) having the chemical composition shown in Table 1, steel plates were fabricated to the thicknesses shown in Tables 2-1 and 2-2 under the manufacturing conditions shown in Tables 2-1 and 2-2. After cold forming, seam welding was performed using submerged arc welding with a welding heat input of approximately 35 to 60 kJ / cm. Subsequently, the pipes were expanded with an expansion ratio of 1.0% to produce steel pipes with an outer diameter of 610 mm and a pipe thickness of 15 to 25 mm. In the manufacturing conditions shown in Tables 2-1 and 2-2, "reduction ratio" refers to the cumulative reduction ratio at temperatures below 900°C, "FT" refers to the rolling end temperature, and "heat treatment" refers to long-term aging treatment. A 1.0% expansion ratio means expanding the inner diameter of the pipe by 1.0% in the diameter direction from the inner surface to the outer surface using a pipe expander. Furthermore, after the cooling stops in the accelerated cooling process, a reheating process is carried out within 150 seconds (s) at the heating rate shown in Tables 2-1 and 2-2.

[0086] As described above, a sample for observing the steel structure was taken from the center of the width of the manufactured steel plate. After mirror polishing the L-section of the steel plate (a section parallel to the rolling direction and parallel to the direction normal to the rolling surface), the crystal orientation of a randomly selected 1 mm × 1 mm region at the center of the plate thickness was measured by electron beam backscatter diffraction (EBSD). The grain size was determined by image analysis, with regions where the angle difference between adjacent pixels was 15° or more being identified as grain boundaries. The EBSD measurement conditions were an acceleration voltage of 17 kV and a measurement pitch of 0.8 μm.

[0087] Note that the average crystal grain size darea (μm) is calculated by comparing it to the area ai (μm) occupied by each crystal grain. 2 ) and the equivalent circular diameter di (μm) of each crystal grain were calculated using the following formula. darea(μm)=Σ(ai·di) / Σai Furthermore, the minimum grain size for the top 20% of bainite grains is calculated by arranging the grains in descending order of their circular diameter (equivalent to the grain size) and selecting 20% ​​of the total grains from largest to smallest, representing the smallest grain size among them.

[0088] Furthermore, after mirror polishing the L-shaped cross-section of the steel plate (a cross-section parallel to the rolling direction and parallel to the direction normal to the rolling surface), the microstructure was revealed by Nital etching. Subsequently, using an optical microscope, a randomly selected 7.1 × 10⁻¹⁴ cross-section was observed at the center of the plate thickness. -2 mm 2 Steel microstructure photographs were taken for five fields of view (magnification: 400x), and the bainite fraction, ferrite fraction, and island martensite fraction in the photographs were measured using an image analysis device (Fiji), and the bainite grain size was measured using an image analysis device (TSL OIM Analysis). The bainite fraction, ferrite fraction, island martensite fraction, and bainite grain size are the average values ​​for the five fields of view. In the aforementioned photographs, regions that are not elongated in the rolling direction and are observed as equiaxed crystal grains were identified as bainite.

[0089] Furthermore, the number of precipitates was determined using the method evaluated in the embodiment, as described above.

[0090] To evaluate the low-temperature toughness of steel plates, DWTT (Deep Water Threshold Test) specimens were taken so that the longitudinal direction of the DWTT specimen was horizontal to the rolling direction and perpendicular to the thickness direction. DWTT was then performed in accordance with API 5L, and the ductile fracture surface ratio and fracture transition temperature at a test temperature of -40°C were evaluated. A good evaluation was given when the ductile fracture surface ratio at -40°C was 85% or higher and the fracture transition temperature was -40°C or lower.

[0091] As a measure of steel plate properties, tensile test specimens were taken so that the longitudinal direction of the tensile test specimen was perpendicular to both the rolling direction and the thickness direction of the steel plate, and tensile tests were performed at 350°C to determine the yield strength and tensile strength. In the tensile test, a round bar specimen with a diameter of 6 mm was used, the crosshead speed was set to 0.15 mm / min, and the yield strength at 350°C was 555 MPa or higher and the tensile strength was 620 MPa. A value of Pa or higher was considered good. The steel plate properties were evaluated by taking test pieces from the steel plate before it was formed into steel pipes. As a characteristic of steel pipes, a 6 mm diameter round bar test specimen was taken with the circumferential direction of the pipe being the longitudinal direction, and a tensile test was performed at 350°C with a crosshead speed of 0.15 mm / min., similar to the steel plate test. Yield strength and tensile strength were evaluated.

[0092] Furthermore, to simulate the high-temperature strength after prolonged holding in the medium 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 where the Larson-Miller Parameter, shown in equation (2), was 15700 (400°C, 2335 hours), which corresponds to holding at 350°C, the applicable temperature for steam piping, for 20 years. Steel plates were evaluated as good if the yield strength at 350°C after heat treatment was 555 MPa or higher and the tensile strength was 620 MPa or higher. The above measurements were performed for both steel plates and steel pipes in the same manner as before heat treatment. The results are shown in Tables 3-1 and 3-2. Furthermore, in order to evaluate the decrease in tensile strength when held at a medium temperature range for a long period of time, the tensile strength of the steel pipe properties was calculated as ((tensile strength before heat treatment (TS0)) - (tensile strength after heat treatment (TS))) / tensile strength before heat treatment (TS0), and a value of 0.050 or less was evaluated as good.

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

[0094] The toughness of the heat-affected zone (HAZ) of the weld was evaluated by Charpy impact testing. The notch position of the Charpy impact test specimen was 3 mm (HAZ 3 mm) from the bond area, which is the boundary between the weld metal and the base metal, towards the base metal. The test temperature was -40°C. In this invention, Charpy impact testing was performed using three test specimens for each condition, and the absorbed energy (vE) at -40°C was determined. -40 A material was evaluated as having excellent toughness if its average value was 60J or higher.

[0095] As described above, Tables 2-1, 2-2, 3-1, and 3-2 show the manufacturing conditions for steel plates as well as the test results for steel plates and steel pipes.

[0096] Examples of the present invention (No. 1 to 34), in which both the component composition and steel sheet manufacturing conditions are within the scope of the present invention, exhibit a yield strength of 555 MPa or higher and a tensile strength of 620 MPa or higher before and after heat treatment (measured at 350°C) of the steel sheet. Furthermore, examples of the present invention (No. 1 to 34) showed good results in terms of low-temperature toughness, HAZ toughness, and (TS0-TS) / TS0 of the steel sheet base material at -40°C. On the other hand, in comparative examples (Nos. 35-67) whose component composition or 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, or (TS0-TS) / TS0 before and after heat treatment (measured at 350°C) of the steel sheet did not reach the target values.

[0097] [Table 1]

[0098] [Table 2-1]

[0099] [Table 2-2]

[0100] [Table 3-1]

[0101] [Table 3-2]

Claims

1. In mass percent, C: 0.040-0.090% 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, The composition has an elemental composition containing N: 0.010% or less, with the remainder being Fe and unavoidable impurities, and the parameter P is represented by the following formula (1). ｅｆｆ The percentage is 0.050% or more, Ti / N is 2.0 to 4.1, and X, represented by formula (3), is 0.69% or more. The microstructure has bainite accounting for 80% or more of the area in the center of the plate 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 is 70 μm or less. The yield strength before and after aging, performed under the condition of Larson Miller Parameter (LMP) = 15700 as defined by equation (2) below, is 555 MPa or higher, and the tensile strength (TS) at 350°C measured after aging and the tensile strength (TS) at 350°C measured before aging are... ０ ) is (TS ０ -TS) / TS ０ The relationship ≤0.050 is satisfied, the toughness of the steel plate is 85% or more as determined by the ductile fracture surface ratio obtained by DWTT at -40°C, and the toughness of the heat-affected zone formed during welding is vE －４０ Steel plates with a temperature of 60J or higher. P ｅｆｆ (%)=(0.13Nb+0.24V-0.125Ti) / (C+0.86N)・・・(1) In equation (1), the element symbols represent the content (mass %) of each element. For elements that are not present, substitute 0. LMP=(T+273)×(20+log(t))...(2) T: Heat treatment temperature (°C) t: Heat treatment time (hours) X=0.35Cr+0.9Mo+12Nb+8V...(3) In equation (3), the element symbols represent the mass percentage of each element present. For elements not present, 0 is substituted.

2. The aforementioned component composition is expressed 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 the following in an amount of 0.0005 to 0.0040%. The steel plate according to claim 1, wherein Y represented by formula (4) is 1.50% or less. Y=Cu+Ni+Cr+Mo...(4) In equation (4), the element symbols represent the content (mass %) of each element. For elements that are not present, 0 is substituted.

3. The microstructure has ferrite with an area ratio of 10% or less and island-like martensite (MA) with an area ratio of 10% or less in 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 from Ti, Nb, V, Mo, Cr, and Al were found in the center of the plate thickness over a 1 mm² area. ２ The steel plate according to claim 1, wherein there are 50,000 to 1,000,000 particles per sheet, 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.

4. The microstructure has ferrite with an area ratio of 10% or less and island-like martensite (MA) with an area ratio of 10% or less in the center of the plate thickness, The steel sheet according to claim 2, wherein 50,000 to 1,000,000 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 from Ti, Nb, V, Mo, Cr, and Al, each with a diameter of 100 nm or less, are present per 1 mm² of the tested area at the center of the sheet thickness, 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.

5. A steel pipe using the steel plate described in any one of claims 1 to 4.

6. A method for manufacturing a steel sheet according to any one of claims 1 to 4, A heating process in which the steel material is heated to 1000-1200°C, A hot rolling step in which the steel material heated in the heating step is hot-rolled under the conditions that the cumulative reduction ratio at 900°C or below is 50% or more, the reduction ratio per pass at 950°C or above is 10% or more in one or more passes, and the reduction ratio per pass at 900°C or below is 15% or more in one or more passes, and the rolling completion temperature is 850°C or below. A method for manufacturing a steel sheet, comprising: an accelerated cooling step in which the hot-rolled steel sheet obtained in the hot-rolling step is accelerated cooling under the conditions that the cooling start temperature is 700°C or higher, the average cooling rate is 10°C / s or higher, and the cooling stop temperature is 250 to 550°C; and a reheating step in which, after cooling has stopped in the accelerated cooling step, the hot-rolled steel sheet is reheated within 150 seconds under the conditions that the heating rate is 1°C / s or higher and the target temperature is 550 to 700°C to obtain a steel sheet.

7. A method for manufacturing a steel pipe, comprising: a cold forming step of cold forming a steel plate according to any one of claims 1 to 4 into a tubular shape; a welding step of welding the butt joints formed by butting the ends of the steel plates formed into a tubular shape in the cold forming step; and a pipe expansion step of expanding the pipe.

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