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

Abstract

The present invention provides a high-strength steel plate that serves as a material for a steel pipe, satisfying all of the following: having a large diameter, having excellent low-temperature toughness, and having strength characteristics required for high-strength steel pipes for steam transport, and a method for manufacturing the same. Furthermore, the present invention provides a steel pipe composed of the high-strength steel plate and a method for manufacturing the same. The high-strength steel plate of the present invention has a specific compositional component and a specific microstructure, and the toughness of the steel plate is 85% or more as the ductile fracture rate obtained by DWTT at -40°C, and the yield strength before and after aging under the condition of LMP = 15700, a Larson Miller parameter defined by the following equation (2), is 555 MPa or more. LMP = (T + 273) × (20 + log(t))···(2) T: Heat treatment temperature (°C) t: Heat treatment time (hours)

Description

Technology Field

[0001] The present invention relates to a high-strength steel plate having excellent low-temperature toughness and a yield strength of 555 MPa or more before and after long-term aging in a medium-temperature range, a method for manufacturing the same, and a steel pipe composed of said high-strength steel plate and a method for manufacturing the same. In particular, the high-strength steel plate of the present invention can be suitably used as a material for high-strength steel pipes for steam piping. Background Technology

[0002] Methods for recovering oil sands from oil layers deposited in Canada and other regions include open-pit mining and steam injection methods, in which high-temperature, high-pressure steam is injected into the oil layer through steel pipes. Since there are few areas where open-pit mining is applicable, steam injection methods are used in many regions. In steam injection methods, the steam injected into the oil layer is in the temperature range of 300 to 400°C (hereinafter referred to as the medium temperature range). In steam injection methods, steam in the medium temperature range is injected into the oil layer at high pressure. As described above, steel pipes are used for the delivery of this steam. Recently, with the aim of improving the recovery rate of heavy oil and reducing installation costs accompanying the increase in energy demand, there is a demand for larger diameters and higher strengths of the above steel pipes. In addition, when laying the above-mentioned steel pipes, laying work and hydrostatic pressure tests are conducted after early spring when temperatures are high; however, excellent low-temperature toughness is also required to enable laying work even during the winter when temperatures are low due to high demand.

[0003] Patent Documents 1 and 2 are prior art for steel pipes for steam transport that can be used in the steam injection method. These patent documents disclose a seamless steel pipe equivalent to API X80 grade, and the outer diameter of this seamless steel pipe is a maximum of 16 inches. In the case of seamless steel pipes, it is difficult to increase the diameter further. In addition, in seamless steel pipes, a large amount of alloying elements is required to obtain strength greater than API X80 grade.

[0004] However, Patent Documents 3 and 4 disclose a manufacturing technology for high-strength steel pipes that can be made large and manufactured by welding. More specifically, Patent Documents 3, 4, and 5 relate to a manufacturing technology for high-strength steel pipes that are produced by a Thermo-mechanical Control Process (TMCP) and have a strength of API X80 grade or higher. Prior art literature

[0005] Japanese Published Patent Application No. 2000-290728 Japanese Patent No. 4821939 Japanese Patent No. 5055736 Japanese Patent No. 4741528 Japanese Patent No. 6241569 The problem to be solved

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

[0007] Patent Document 4 describes the creep properties of steel pipes after long-term aging, and while the developed steel shows improvement in fracture strength, the criteria for judging safety are not clearly established. Furthermore, in several places, the fracture strength is found to be lower than 80% (=440 MPa) of the lower limit of the yield stress specification for API X80 grade, so it is difficult to say that it is sufficient strength.

[0008] In the steel pipe disclosed in Patent Document 5, although the high-temperature characteristics in the medium temperature range satisfy the API X80 grade, the LMP is small, and the strength characteristics when the steel pipe is used for a long time are not considered, and the low-temperature toughness is not considered either.

[0009] As such, in conventional technology, it is not possible to obtain a high-strength steel pipe for steam piping that satisfies all of the following: being large in diameter, having strength characteristics required for high-strength steel pipes for steam transport, and having excellent low-temperature toughness.

[0010] The present invention aims to solve the above problem by providing a high-strength steel plate that serves as a material for a steel pipe, satisfying all of the following: having a large diameter, having excellent low-temperature toughness, and having strength characteristics required for high-strength steel pipes for steam transport, and a method for manufacturing the same. In addition, the present invention also aims to provide a steel pipe composed of the high-strength steel plate and a method for manufacturing the same.

[0011] In addition, "excellent low-temperature toughness" means that the ductile fracture rate obtained by DWTT (test temperature: -40°C) in accordance with API 5L is 85% or higher, and the fracture transition temperature is -40°C or lower. The test temperature for DWTT was set to -40°C to account for the decrease in toughness due to work hardening during tube forming. means of solving the problem

[0012] The inventors carefully examined the characteristics of high-strength steel sheets for large-diameter steel pipes in the medium temperature range. As a result, they discovered that by appropriately selecting the compositional composition and manufacturing conditions, a high-strength steel sheet capable of producing a high-strength steel pipe with a large diameter and possessing the strength characteristics and low-temperature toughness required for high-strength steel pipes for steam transport can be obtained.

[0013] The present invention is formed by adding further consideration to the above recognition, that is, the present invention is composed as follows.

[0014] [1] The composition of the components is in mass %,

[0015] C: 0.04∼0.09%,

[0016] Si: 0.03∼0.25%,

[0017] Mn: 1.5∼2.5%,

[0018] P: 0.020% or less,

[0019] S: 0.002% or less,

[0020] Mo: 0.10∼0.50%,

[0021] Nb: 0.010∼0.055%,

[0022] Ti: 0.005∼0.020%,

[0023] Ca: 0.0040% or less,

[0024] Al: 0.01∼0.04%,

[0025] N: Contains 0.006% or less, and the remainder consists of Fe and unavoidable impurities,

[0026] X(%) represented by Equation (1) is 0.65% or more, and

[0027] The microstructure has bainite with an area ratio of 80% or more at the center of the plate thickness, the average crystal grain size of the bainite is 25㎛ or less, and the minimum crystal grain size of the top 20% of the crystal grains with large grain sizes among the bainite is 60㎛ or less, and the toughness of the steel plate has a ductile fracture surface rate of 85% or more obtained by DWTT at -40℃, and

[0028] A high-strength steel plate having a yield strength of 555 MPa or more before and after aging under the condition of LMP = 15700, a Larson Miller parameter defined by the following (2).

[0029] X(%)=0.35Cr+0.9Mo+12.5Nb+8V···(1)

[0030] In Equation (1), the element symbols represent the content (mass%) of each element. Additionally, 0 is substituted for elements that are not contained.

[0031] LMP=(T+273)×(20+log(t))···(2)

[0032] T: Heat treatment temperature (°C)

[0033] t: Heat treatment time (time)

[0034] [2] The above composition of components, in mass%, additionally,

[0035] Cr: 0.50% or less,

[0036] V: 0.070% or less,

[0037] Cu: 0.50% or less,

[0038] A high-strength steel plate as described in [1] containing one or more of the following: 0.50% or less of Ni.

[0039] [3] Additionally, the plate has ferrite with an area ratio of 10% or less in the center of the plate thickness, and island-shaped martensite with an area ratio of 10% or less,

[0040] At the center of the plate thickness, the total number of Ti-based precipitates with a diameter of 100 nm or less, Nb-based precipitates with a diameter of 100 nm or less, V-based precipitates with a diameter of 100 nm or less, Mo-based precipitates with a diameter of 100 nm or less, Cr-based precipitates with a diameter of 100 nm or less, Al-based precipitates with a diameter of 100 nm or less, and composite precipitates with a diameter of 100 nm or less containing two or more elements among Ti, Nb, V, Mo, Cr, and Al is 50,000 to 1,000,000 per 1 mm² of observed area, and

[0041] A high-strength steel plate as described in claim 1 or 2, wherein 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.

[0042] [4] A steel pipe using a high-strength steel plate as described in any of [1] to [3] above.

[0043] [5] A method for manufacturing a high-strength steel plate as described in any of [1] to [3] above, comprising a heating process of heating a steel material to 1000 to 1200°C, and

[0044] A hot rolling process for hot rolling a steel material heated in the above heating process, comprising at least one rolling pass at 950°C or higher in which the reduction rate per pass is 10% or more, and at least one rolling pass at 900°C or lower in which the reduction rate per pass is 15% or more, and also a cumulative reduction rate at 900°C or lower in which the rolling end temperature is 50% or more, and a rolling end temperature of Ar3 or higher in which 850°C or lower;

[0045] A method for manufacturing a high-strength steel sheet having an accelerated cooling process in which a hot-rolled steel sheet obtained from the above hot rolling process is accelerated under conditions in which the cooling start temperature is Ar3 or higher, the average cooling rate is 5℃ / s or higher, and the cooling stop temperature is 300 to 550℃.

[0046] [6] A cold forming process for cold forming a high-strength steel plate into a tubular shape as described in any of [1] to [3] above, and

[0047] A welding process for welding the butt joints formed by butting the ends of steel plates formed into a tubular shape in the above cold forming process.

[0048] A method for manufacturing a steel pipe having Effects of the invention

[0049] According to the present invention, a high-strength steel pipe is obtained that has a large diameter and possesses the strength and toughness characteristics required for high-strength steel pipes for steam transport. Specific details for implementing the invention

[0050] (Form for carrying out the invention)

[0051] Embodiments of the present invention are described below. Furthermore, the present invention is not limited to the following embodiments.

[0052] High-strength steel plate

[0053] The high-strength steel sheet of the present invention (also simply referred to as a steel sheet) contains, in mass%, C: 0.04 to 0.09%, Si: 0.03 to 0.25%, Mn: 1.5 to 2.5%, P: 0.020% or less, S: 0.002% or less, Mo: 0.10 to 0.50%, Nb: 0.010 to 0.055%, Ti: 0.005 to 0.020%, Ca: 0.0040% or less, Al: 0.01 to 0.04%, and N: 0.006% or less. In the following description, "%" indicating the content of a component means "mass%".

[0054] C: 0.04∼0.09%

[0055] Carbon (C) is an element necessary to secure the strength of steel through solid solution strengthening and precipitation strengthening. In particular, an increase in the amount of solid solution carbon and the formation of precipitates contribute to securing strength in the medium temperature range. In order to secure a predetermined strength at room temperature and in the medium temperature range, the present invention sets the carbon content to 0.04% or more. It is preferable that the carbon content be 0.05% or more. Furthermore, if the carbon content exceeds 0.09%, it results in a deterioration of toughness and weldability. For this reason, the upper limit of the carbon content is set to 0.09%. Additionally, it is preferable that the carbon content be 0.08% or less, and more preferable that it be 0.07% or less.

[0056] Si: 0.03∼0.25%

[0057] Si is added for deoxidation. If the Si content is less than 0.03%, a sufficient deoxidation effect is not obtained, and the toughness of the steel sheet decreases. For this reason, the Si content is set to 0.03% or more. It is preferable for the Si content to be 0.04% or more, more preferable for 0.05% or more, even more preferable for 0.07% or more, and most preferable for 0.10% or more. On the other hand, if the Si content exceeds 0.25%, it causes a deterioration in toughness. Therefore, the Si content is set to 0.25% or less. Furthermore, it is preferable for the Si content to be 0.23% or less, more preferable for 0.20% or less, even more preferable for 0.19% or less, and most preferable for 0.18% or less.

[0058] Mn: 1.5∼2.5%

[0059] Manganese is an element effective for improving the strength and toughness of steel. The effect is small when the Mn content is less than 1.5%. For this reason, the Mn content is set to 1.5% or more. It is preferable to have a Mn content of 1.55% or more, more preferable to 1.6% or more, even more preferable to 1.65% or more, and most preferable to 1.7% or more. On the other hand, if the Mn content exceeds 2.5%, toughness and weldability deteriorate significantly. Therefore, the Mn content is set to 2.5% or less. Furthermore, it is preferable to have a Mn content of 2.4% or less, more preferable to 2.3% or less, even more preferable to 2.2% or less, and most preferable to 2.1% or less.

[0060] P: 0.020% or less

[0061] P is an impurity element that degrades toughness. For this reason, 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 where toughness deterioration falls within an acceptable range, the P content is set to 0.020% or less. It is preferable to set the P content to 0.018% or less, more preferable to 0.015% or less, even more preferable to 0.013% or less, and most preferable to 0.010% or less. Furthermore, regarding the lower limit of the P content, it is not particularly limited, and 0% is acceptable, or it is more preferable to set it to 0.003% or more.

[0062] S: 0.002% or less

[0063] S is an impurity element that degrades toughness. For this reason, it is desirable to reduce the S content as much as possible. Furthermore, even if shape control is performed by adding Ca to convert MnS into CaS-based inclusions, in the case of X80 grade high-strength materials, finely dispersed CaS-based inclusions can also be a factor in the reduction of toughness. For this reason, the S content is kept at 0.002% or less. It is desirable to keep the S content at 0.0018% or less, more desirable to keep it at 0.0015% or less, even more desirable to keep it at 0.0013% or less, and most desirable to keep it at 0.0010% or less. Additionally, there is no particular limitation on the lower limit of the S content; it may be 0%, or it is desirable to have 0.0003% or more.

[0064] Mo: 0.10∼0.50%

[0065] Mo strengthens steel through solid solution and also improves hardenability. Through solid solution strengthening and improved hardenability, an increase in strength is achieved, and in particular, strength in the medium temperature range increases due to increased resistance to tempering softening. When the Mo content is less than 0.10%, the effect is small, and sufficient strength is not obtained. For this reason, the Mo content is set to 0.10% or more. It is preferable to set the Mo content to 0.13% or more, more preferable to 0.15% or more, even more preferable to 0.18% or more, and most preferable to 0.20% or more. On the other hand, if the Mo content exceeds 0.50%, the effect becomes saturated, and toughness and weldability deteriorate. Therefore, the Mo content is set to 0.50% or less. In addition, the Mo content is preferably 0.45% or less, more preferably 0.40% or less, even more preferably 0.35% or less, and most preferably 0.30% or less.

[0066] Nb: 0.010∼0.055%

[0067] Nb inhibits grain growth during slab heating and rolling. Consequently, the microstructure is refined, allowing sufficient strength and toughness to be imparted to the steel. Additionally, Nb is a necessary component for securing strength in the medium temperature range by forming carbides. This effect is significant when the Nb content is 0.010% or higher. For this reason, the Nb content is set to 0.010% or higher. It is preferable that the Nb content be 0.013% or higher, more preferable that it be 0.015% or higher, even more preferable that it be 0.018% or higher, and most preferable that it be 0.020% or higher. On the other hand, if the Nb content exceeds 0.055%, not only is the effect nearly saturated, but toughness and weldability also deteriorate. Therefore, the Nb content is set to 0.055% or lower. It is preferable that the Nb content be 0.053% or less, more preferable that it be 0.050% or less, even more preferable that it be 0.048% or less, and most preferable that it be 0.045% or less.

[0068] Ti: 0.005∼0.020%

[0069] Ti forms TiN, which suppresses grain growth during slab heating or in the weld heat-affected zone, thereby refining the microstructure and improving toughness. To obtain this effect, it is necessary to have a Ti content of 0.005% or more. It is preferable for the Ti content to be 0.006% or more, more preferable for 0.007% or more, even more preferable for 0.008% or more, and most preferable for 0.009% or more. On the other hand, if the Ti content exceeds 0.020%, toughness deteriorates. Therefore, the Ti content is kept at 0.020% or less. It is preferable for the Ti content to be 0.019% or less, more preferable for 0.018% or less, even more preferable for 0.017% or less, and most preferable for 0.016% or less.

[0070] Ca: 0.0040% or less

[0071] Ca has the effect of improving the toughness of steel sheets by fixing sulfur in the steel. To obtain this effect, it is desirable to contain 0.0010% or more of Ca, and more desirable to contain 0.0015% or more. However, if the Ca content exceeds 0.0040%, the inclusions in the steel increase, and toughness may deteriorate. Therefore, the Ca content is kept at 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.

[0072] Al: 0.01∼0.04%

[0073] Al is added as a deoxidizing agent. If the Al content is less than 0.01%, a sufficient deoxidation effect is not obtained, and the toughness of the steel sheet decreases. The Al content should be 0.01% or more, preferably 0.013% or more, more preferably 0.015% or more, even more preferably 0.018% or more, and most preferably 0.02% or more. On the other hand, if the Al content exceeds 0.04%, the toughness deteriorates. Therefore, the Al content should be 0.04% or less. The Al content should be 0.038% or less, more preferably 0.035% or less, even more preferably 0.032% or less, and most preferably 0.03% or less.

[0074] N: 0.006% or less

[0075] N forms TiN together with Ti and is finely dispersed in the high-temperature region of the weld heat-affected zone reaching 1350°C or higher. This fine dispersion refines the austenite grains in the weld heat-affected zone, significantly contributing to the improvement of toughness in the weld heat-affected zone. To obtain this effect, it is desirable for the N content to be 0.002% or higher, and more desirable to be 0.0025% or higher. On the other hand, if the N content exceeds 0.006%, it leads to coarsening of precipitates and deterioration of base metal toughness due to an increase in dissolved N, as well as deterioration of weld metal toughness in steel pipes. Therefore, the N content is kept at 0.006% or less. It is desirable for the N content to be 0.0058% or less, more desirable to be 0.0055% or less, even more desirable to be 0.0053% or less, and most desirable to be 0.005% or less.

[0076] X(X(%)=0.35Cr+0.9Mo+12.5Nb+8V···(1)): 0.65% or more

[0077] In the present invention, the Cr content, Mo content, Nb content, and V content are adjusted so that X(%), represented by the following formula (1), is 0.65% or more. X(%) is an important factor for making the steel have excellent strength after long-term aging in the medium temperature range. Therefore, it is necessary to make X(%) 0.65% or more in order to suppress dislocation recovery during long-term aging. It is preferable to make X(%) 0.66% or more, more preferable to make it 0.67% or more, even more preferable to make it 0.68% or more, and most preferable to make it 0.69% or more. In order to manufacture the high-strength steel sheet of the present invention at a low cost, it is preferable to make X(%) 1.20% or less, more preferable to make it 1.19% or less, even more preferably 1.18% or less, and most preferably 1.17% or less. In addition, there are cases in which Cr and V are not included in the present invention, but in such cases, it is acceptable to substitute 0 for “Cr” or “V” in Equation (1).

[0078] X(%)=0.35Cr+0.9Mo+12.5Nb+8V···(1)

[0079] In Equation (1), the element symbols represent the content (mass%) of each element. Additionally, 0 is substituted for elements that are not contained.

[0080] In addition, the high-strength steel plate of the present invention may contain one or more of Cr, V, Cu, and Ni for the purpose of further improving properties.

[0081] Cr: 0.50% or less

[0082] Cr is one of the elements effective for increasing high-temperature strength by increasing resistance to tempering softening. On the other hand, 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. It is preferable that the Cr content be 0.48% or less, more preferable that it be 0.45% or less, even more preferable that it be 0.43% or less, and most preferable that it be 0.40% or less. Furthermore, if the resistance to tempering softening and high-temperature strength can be increased by X being within the desired range, the lower limit of the Cr content is not particularly limited and 0% is acceptable. To obtain the above effect, it is preferable that the Cr content be 0.05% or more, more preferable that it be 0.08% or more, and even more preferable that it be 0.10% or more.

[0083] V: 0.070% or less

[0084] Adding a small amount of V refines the grain size, contributing to an increase in strength. Additionally, it increases resistance to tempering softening, making it one of the elements effective for increasing strength in the medium temperature range. On the other hand, if the V content exceeds 0.070%, the toughness of the weld heat-affected zone deteriorates. Therefore, when V is included, the V content should be 0.070% or less. The V content is preferably 0.060% or less, more preferably 0.050% or less, even more preferably 0.040% or less, and most preferably 0.030% or less. Furthermore, if high-temperature strength can be increased by having X within a desired range, the lower limit of the V content is not particularly limited and can be 0%. To obtain the above effects, the V content is preferably 0.005% or more, and more preferably 0.010% or more.

[0085] Cu: 0.50% or less

[0086] Cu is one of the elements effective for improving toughness and increasing strength. To obtain these effects, it is desirable to have a Cu content of 0.05% or more, and more desirable to have a Cu content of 0.10% or more. On the other hand, if the Cu content exceeds 0.50%, weldability decreases. Therefore, when Cu is included, the Cu content should be 0.50% or less. It is desirable to have a Cu content of 0.48% or less, more desirable to have a Cu content of 0.45% or less, even more desirable to have a Cu content of 0.43% or less, and most desirable to have a Cu content of 0.40% or less.

[0087] Ni: 0.50% or less

[0088] Ni is one of the elements effective for improving toughness and increasing strength. To obtain these effects, it is desirable to have a Ni content of 0.05% or more, and more desirable to have a Ni content of 0.10% or more. If the Ni content exceeds 0.50%, the effect becomes saturated, leading to an increase in manufacturing costs. Therefore, when Ni is included, the Ni content is kept at 0.50% or less. It is desirable to have a Ni content of 0.48% or less, more desirable to have a Ni content of 0.45% or less, even more desirable to have a Ni content of 0.43% or less, and most desirable to have a Ni content of 0.40% or less.

[0089] In the present invention, it is preferable that the Cu content, Ni content, Cr content, and Mo content be adjusted so that Cu + Ni + Cr + Mo (element symbols indicate the content (mass%) of each element) is 1.00% or less. In order to achieve a balance between the effect of strength increase by these element groups and manufacturing costs, it is preferable to keep Cu + Ni + Cr + Mo at 1.00% or less, more preferable to keep it at 0.98% or less, even more preferable to keep it at 0.95% or less, and most preferable to keep it at 0.90% or less. Furthermore, although there is no particular limitation regarding the lower limit, it is preferable to keep Cu + Ni + Cr + Mo at 0.10% or more, more preferable to keep it at 0.15% or more, even more preferable to keep it at 0.20% or more, and most preferable to keep it at 0.25% or more.

[0090] The remainder other than the above components is Fe and unavoidable impurities. These are impurities inevitably incorporated from raw materials, manufacturing processes, or manufacturing equipment, and are permitted to be included within a range that does not impede the purpose 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, REM, etc.

[0091] Next, the structure of the high-strength steel plate of the present invention will be described.

[0092] Bainite is 80% or more in area ratio at the center of the plate thickness

[0093] Bainite is an important microstructure for achieving both strength and low-temperature toughness. Furthermore, bainite effectively contributes to the improvement of steel sheet strength through the reinforcement of the transformation microstructure. Additionally, uniformity of the microstructure is required to increase the initial dislocation density and to improve the strength of high-strength steel, particularly in the medium-temperature range. For this reason, the microstructure of the high-strength steel of the present invention must be bainite-dominated; specifically, the bainite must account for at least 80% of the total steel microstructure in terms of area percentage at the center of the sheet thickness. It is preferable that the bainite account for at least 83% of the area percentage, more preferable that it account for at least 85%, even more preferable that it account for at least 88%, and most preferable that it account for at least 90%. Although there is no specific upper limit on the bainite fraction, from the perspective of improving deformation performance, it is preferable that the bainite account for at least 98% of the area percentage, and more preferable that it account for at least 95%.

[0094] Ferrite: Area percentage 10% or less (suitable condition)

[0095] When ferrite is processed into processed ferrite through rolling, the number of mobile dislocations increases, which can improve strength. On the other hand, since the dislocation density of the ferrite phase significantly decreases due to heat treatment in the medium temperature range, the ferrite phase becomes softer, which may lead to a decrease in the strength of the steel sheet after heat treatment. Therefore, it is desirable to keep the ferrite area percentage in the center of the sheet thickness 10% or less. The area percentage of ferrite is more preferably 9% or less, even more preferably 8% or less, and most preferably 7% or less. Although there is no specific lower limit for the area percentage of ferrite, from the perspective of improving deformation performance, it is desirable for the ferrite area percentage to be 1% or more.

[0096] Island-shaped martensite (MA): Area percentage 10% or less (suitable condition)

[0097] Martensite-austenite constituents (MA) are very hard phases and may reduce the low-temperature toughness of steel sheets by acting as fracture initiation points. Therefore, it is desirable to keep the area percentage of martensite (MA) in the center of the sheet thickness 10% or less. More preferably, the area percentage of martensite is 8% or less, even more preferably 5% or less, and most preferably 3% or less. Although there is no specific lower limit for the area percentage of martensite, it is desirable to have an area percentage of 1% or more, and more desirable to have an area percentage of 2% or more.

[0098] In addition, as mentioned above, the steel structure of the steel plate that serves as the base material basically needs to be composed of the bainite mentioned above, but as a remainder structure other than bainite, ferrite, and island-shaped martensite (MA), pearlite, martensite, cementite, and retained austenite may be cited.

[0099] The average crystal grain size of bainite is 25㎛ or less, and the minimum crystal grain size of the top 20% of crystal grains with the largest grain sizes among the bainite is 60㎛ or less.

[0100] Since the grain boundaries of bainite act as a resistance to brittle crack propagation, grain refinement contributes to the improvement of low-temperature toughness. Therefore, the average grain size of all bainites is set to 25 μm or less. It is preferable that the average grain size of all bainites be 23 μm or less, more preferable that it be 20 μm or less, even more preferable that it be 19 μm or less, and most preferable that it be 18 μm or less. Regarding the lower limit, there is no particular limitation, but it is preferable that it be 5 μm or more, and more preferable that it be 6 μm or more.

[0101] In addition, while low-temperature toughness is improved by refining the average crystal grain size, there is a limit to the refining of the average crystal grain size when cooling is initiated near the Ar3 temperature. In the present invention, it is important to further suppress the formation of coarse crystal grains. Coarse bainite is highly likely to become a fracture initiation point, and if the crystal grain size of the top 20% with the largest grain size among all bainites is large, low-temperature toughness deteriorates. In particular, since it is easy to become a fracture initiation point when the minimum crystal grain size of the top 20% with the largest grain size exceeds 60 μm, it is necessary to make the minimum crystal grain size of the top 20% with the largest grain size at the center of the plate thickness 60 μm or less, preferably 55 μm or less, more preferably 50 μm or less, even more preferably 45 μm or less, and most preferably 40 μm or less. Regarding the lower limit, there is no particular limitation, but 10 μm or more is preferred, and 15 μm or more is more preferred.

[0102] Here, the average crystal grain size of bainite was determined as follows.

[0103] That is, after mirror polishing the L cross-section of the steel plate (a cross-section parallel to the rolling direction and parallel to the normal direction of the rolling plane), the crystal orientation of a randomly selected 1 mm × 1 mm² area at the center of the plate thickness was measured by electron backscatter diffraction (EBSD), and the result was obtained by image analysis using the area where the angle difference between adjacent pixels is 15° or more as the crystal grain boundary. The EBSD measurement conditions were set to an acceleration voltage of 17 kV and a measurement pitch of 0.8 μm.

[0104] In addition, the average crystal grain size darea (㎛) is the area ai (㎛) occupied by each crystal grain. 2 It was calculated from the original diameter di (㎛) of each crystal grain and the following formula.

[0105] darea(㎛)=Σ(ai·di) / Σai

[0106] In addition, the minimum crystal grain size in the top 20% of crystal grains with large grain sizes in bainite represents the minimum crystal grain size when the equivalent diameters of each crystal grain are arranged in order of largest to smallest, and 20% of the total number of crystal grains is selected from the largest.

[0107] The total number of Ti-based precipitates with a diameter of 100 nm or less, Nb-based precipitates with a diameter of 100 nm or less, V-based precipitates with a diameter of 100 nm or less, Mo-based precipitates with a diameter of 100 nm or less, Cr-based precipitates with a diameter of 100 nm or less, Al-based precipitates with a diameter of 100 nm or less, and complex precipitates with a diameter of 100 nm or less composed of two or more elements among Ti, Nb, V, Mo, Cr, and Al, is 50,000 to 1,000,000 per 1 mm² of the inspection area (suitable condition)

[0108] The reasons for limiting the size and total number of Ti-based precipitates, Nb-based precipitates, V-based precipitates, Mo-based precipitates, Cr-based precipitates, Al-based precipitates, and composite precipitates containing two or more elements among Ti, Nb, V, Mo, Cr, and Al (e.g., 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 are explained.

[0109] Even if precipitates (fine precipitates) with a diameter of 100 nm or less are formed, if the number of such precipitates is less than 50,000 per 1 mm², the effect of inhibiting austenite grain growth is weak, and there may be cases where the decrease in strength in the medium temperature range cannot be suppressed. Therefore, it is desirable to have at least 50,000 fine precipitates with a diameter of 100 nm or less in the steel at the center of the plate thickness. More preferably, at least 80,000, even more preferably at least 100,000, and most preferably at least 130,000. On the other hand, if the number of precipitates with a diameter of 100 nm or less exceeds 1,000,000 per 1 mm², the fine precipitates aggregate and coarsen, which conversely weakens the effect of inhibiting austenite grain growth, and there may be cases where the decrease in strength in the medium temperature range becomes greater. Accordingly, it is desirable to have no more than 1,000,000 fine precipitates with a diameter of 100 nm or less per 1 mm². More preferably, no more than 950,000, even more preferably no more than 900,000, and most preferably no more than 850,000.

[0110] Here, the method for measuring the density and size (diameter) of precipitates is described. For the etched surface at any location on the head cross-section, observation is performed using a scanning electron microscope (SEM), or by preparing extraction replica samples or thin film samples and observing them using a transmission electron microscope (TEM). The number of precipitates smaller than 100 nm is determined to be at least 100 µm 2 Measurements are taken within the above area. For example, when observing with a field of view of 2,000 nm × 2,000 nm at a magnification of 100,000 times, the observation area per field of view is 4 µm. 2 Therefore, 25 fields of view are observed randomly. These measurement results are converted into the number per unit area. The number of precipitates smaller than 100 nm is 25 fields of view (100 µm 2If there are 5 of them, the density of the precipitates can be converted to 50,000 per 1 mm². The density of the precipitates is referred to as the number of precipitates per 1 mm² of the inspection area. In addition, the diameter of the precipitates is defined as the average value of the major axis and the minor axis.

[0111] Next, the strength characteristics of the high-strength steel sheet of the present invention will be described. The high-strength steel sheet of the present invention has a yield strength of 555 MPa or more before and after aging under the condition of LMP = 15700, which is the Larson Miller Parameter.

[0112] The aging treatment of LMP = 15700 is an aging treatment performed under the conditions of a heat treatment temperature and heat treatment time such that the LMP represented by the following equation (2) becomes 15700. LMP = 15700 corresponds to the condition of heat treatment for 20 years at 350°C, which is the medium temperature range. In addition, the condition of Larson Miller Parameter (LMP) = 15700 refers to the Larson Miller Parameter (LMP) being 15650 or higher and less than 15750.

[0113] LMP=(T+273)×(20+log(t))···(2)

[0114] T: Heat treatment temperature (°C)

[0115] t: Heat treatment time (time)

[0116] The high-strength steel plate of the present invention has a yield strength of 555 MPa or higher before and after the aging treatment. Since the yield strength is 555 MPa or higher, it has the effect of enabling stable operation as a steel pipe for steam piping. Here, the yield strength refers to the yield strength measured in a high-temperature tensile test at 350°C. The yield strength before and after the aging treatment is preferably 560 MPa or higher, more preferably 565 MPa or higher, even more preferably 570 MPa or higher, and most preferably 575 MPa or higher. Although the upper limit is not particularly limited, the yield strength is preferably 840 MPa or lower.

[0117] In addition, as described in the examples, in the present invention, the yield strength of both the test specimen taken from the steel plate and the steel pipe before and after the aging treatment is 555 MPa or higher.

[0118] 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 (compliance requirement)

[0119] In the present invention, it is desirable 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 indicator for evaluating the decrease in yield strength when maintained in the medium temperature range for a long time. If this difference is 50 MPa or less, the decrease in yield strength after maintaining in the medium temperature range for a long time is within a range where there are no practical problems. If this difference is greater than 50 MPa, a significant decrease in yield strength occurs after maintaining in the medium temperature range for a long time. For this reason, it is desirable to keep the difference in yield strength before and after aging 50 MPa or less. The difference in yield strength before and after aging is more preferably 45 MPa or less, even more preferably 40 MPa or less, and most preferably 35 MPa or less. The lower limit is not specifically limited and can be a negative value or greater than -100 MPa.

[0120] In addition, as an aging treatment condition under the above conditions, for example, heat treatment at 400°C for 2335 hours can be cited.

[0121] The toughness of the steel sheet is a ductile fracture rate of 85% or more obtained by DWTT at -40℃.

[0122] The toughness of the steel sheet of the present invention is such that the ductile fracture rate obtained by DWTT (test temperature: -40°C) in accordance with API 5L is 85% or higher. If the above ductile fracture rate is less than 85%, the steel sheet is prone to brittle fracture at low temperatures, making it difficult to perform laying work throughout the year, including winter conditions where the temperature falls below 0°C, or to use it in regions with environments where the ambient temperature is very low; therefore, it is necessary to ensure the ductile fracture rate is 85% or higher. Furthermore, a ductile fracture rate of 85% or higher obtained by DWTT at -40°C means that the fracture transition temperature is -40°C or lower. It is preferable that the ductile fracture rate be 86% or higher, more preferable that it be 87% or higher, even more preferable that it be 88% or higher, and most preferable that it be 89% or higher. In the above DWTT, the test temperature was set to -40°C to account for the anticipated decrease in toughness due to work hardening during tube forming. The upper limit of the ductile fracture rate obtained by DWTT at -40°C is not specifically limited and may be 100% or less.

[0123] Steel pipe

[0124] Since the steel pipe of the present invention is manufactured using the high-strength steel plate of the present invention, it possesses the strength characteristics and low-temperature toughness required for high-strength welded steel pipes for steam transport, even in large diameters.

[0125] Large diameter means that the outer diameter (diameter) of the steel pipe is 400 mm or more. It is preferable that the outer diameter of the steel pipe be 500 mm or more, more preferable that it be 600 mm or more, and even more preferable that it be 700 mm or more. Although the maximum outer diameter is not particularly limited, it may be 1500 mm or less, and more preferable that it be 1400 mm or less. According to the present invention, it is possible to increase the diameter while maintaining the strength characteristics required for high-strength welded steel pipes for steam transport.

[0126] In addition, the thickness of the steel pipe is not particularly limited, but for steam transport, it is 12 to 30 mm. That is, the thickness of the steel pipe is preferably 12 mm or more, more preferably 13 mm or more, even more preferably 14 mm or more, and most preferably 15 mm or more. In addition, 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.

[0127] The strength characteristics required for high-strength welded steel pipes for steam transport are such that, similar to the high-strength steel, the yield strength at 350°C before and after aging under the condition of Larson Miller Parameter (LMP) = 15700 is 555 MPa or higher. It is preferable that it be 560 MPa or higher, more preferable that it be 565 MPa or higher, even more preferable that it be 570 MPa or higher, and most preferable that it be 575 MPa or higher.

[0128] Although the upper limit is not specifically limited, the yield strength at 350°C before and after the above aging period is preferably 840 MPa or less, and more preferably 830 MPa or less.

[0129] Method for manufacturing high-strength steel plates

[0130] Next, the method for manufacturing a high-strength steel sheet according to the present invention will be described. The method for manufacturing a high-strength steel sheet according to the present invention comprises a heating process, a hot rolling process, and an accelerated cooling process. Each process will be described below. Furthermore, in the following description, unless otherwise specifically stated, the temperature is the average temperature in the thickness direction. The average temperature in the thickness direction is obtained from the surface temperature of the slab or steel sheet by simulation calculations, etc., using parameters such as plate thickness, cooling conditions, and heat transfer rate. For example, the average temperature in the thickness direction is obtained by calculating the temperature distribution in the thickness direction using the difference method. The average temperature in the thickness direction is the average value of each temperature in the thickness direction from the steel sheet surface to the steel sheet surface on the opposite side. In addition, the cooling rate is the average cooling rate obtained by dividing the temperature difference required to cool down to the cooling stop (termination) temperature after the end of hot rolling by the time required to perform such cooling.

[0131] Heating process

[0132] In the method for manufacturing a high-strength steel plate according to 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 becomes the composition of the high-strength steel plate, the adjustment of the composition of the high-strength steel plate may be performed at the stage of adjusting the composition of the molten steel. Furthermore, the steelmaking method of the steel material is not particularly limited.

[0133] In the hot rolling process described below, in order to sufficiently promote austenitization and solid solution of carbides to obtain sufficient strength in the room temperature and medium temperature ranges, it is necessary to heat the steel material to 1000°C or higher. The heating temperature is preferably 1010°C or higher, more preferably 1020°C or higher, even more preferably 1030°C or higher, and most preferably 1040°C or higher. On the other hand, if the heating temperature exceeds 1200°C, austenite grain growth becomes significant, and the toughness of the base material deteriorates. Therefore, the heating temperature is set to 1200°C or lower. Furthermore, the heating temperature is preferably 1190°C or lower, more preferably 1180°C or lower, even more preferably 1170°C or lower, and most preferably 1160°C or lower.

[0134] Hot rolling process

[0135] In the method for manufacturing a high-strength steel sheet according to the present invention, the hot rolling process is a process of hot rolling a steel material heated in a heating process, wherein the rolling includes at least one pass of rolling with a reduction rate of 10% or more per pass at 950°C or higher and at least one pass of rolling with a reduction rate of 15% or more per pass at 900°C or lower, and also includes a cumulative reduction rate of 50% or more at 900°C or lower and a rolling end temperature of 850°C or higher at Ar3 temperature or higher.

[0136] The upper limit of the austenite unrecrystallization temperature range rises to around 900°C due to the Nb content. By performing rolling in the temperature range below 900°C, which is the austenite unrecrystallization range, the austenite grains elongate and become finer in the thickness and width directions, and the density of dislocations within the grains introduced by rolling increases. Therefore, the cumulative reduction rate below 900°C is set to 50% or more. It is preferable that the cumulative reduction rate below 900°C be 55% or more, more preferable that it be 60% or more, even more preferable that it be 65% or more, and most preferable that it be 70% or more. Although the upper limit of the cumulative reduction rate is not particularly limited, it is preferable that it be 95% or less, more preferable that it be 90% or less, and even more preferable that it be 85% or less.

[0137] In addition, this effect is significantly exhibited by having a cumulative reduction rate of 50% or more at 900℃ or lower, and by setting the rolling end temperature to 850℃ or higher than the Ar3 temperature. Furthermore, in high-strength steel sheets manufactured by hot rolling and the cooling described below, and in steel pipes made of such high-strength steel sheets, strength, particularly strength in the medium temperature range, is improved. In addition, the Ar3 temperature can be calculated by the following formula.

[0138] Ar3(℃)=910-310[C]-80[Mn]-20[Cu]-55[Ni]-15[Cr]-80[Mo]

[0139] However, [C], [Mn], [Cu], [Ni], [Cr], and [Mo] in the formula each represent the content (mass%) of C, Mn, Cu, Ni, Cr, and Mo in the base steel sheet (high-strength steel sheet). Additionally, if some elements are not contained in the base steel sheet, the content of said element may be set to "0" to determine the Ar3 temperature.

[0140] In addition, the upper limit of the above cumulative reduction rate is not specifically limited, but since there is a risk that an excessive load will be applied to the rolling mill if the cumulative reduction rate is excessively large, it is desirable to keep the cumulative reduction rate at 90% or less.

[0141] In addition, by including at least one rolling pass in which the reduction rate per pass at 950°C or higher is 10% or more, the fineness of austenite grains due to recrystallization proceeds, thereby improving the low-temperature toughness of the steel. Therefore, it is necessary to include at least one rolling pass in which the reduction rate per pass at 950°C or higher is 10% or more. Furthermore, it is preferable to include at least two rolling passes in which the reduction rate per pass at 950°C or higher is 10% or more. It is more preferable to include at least three rolling passes in which the reduction rate per pass at 950°C or higher is 10% or more, and even more preferable to include at least four rolling passes. Additionally, while there is no specific upper limit on the number of rolling passes in which the reduction rate per pass is 10% or more, it may be 50 passes or less, it is more preferable to be 30 passes or less, and even more preferable to be 15 passes or less. In addition, the reduction rate per pass at 950°C or higher is not specifically limited to an upper limit, but it is preferable that it be 40% or less.

[0142] In addition, by including at least one rolling pass in which the reduction rate per pass at 900°C or lower is 15% or more, deformation is introduced into the grains, thereby promoting the fineness of the austenite grains and improving the low-temperature toughness of the steel. Therefore, it is necessary to include at least one rolling pass in which the reduction rate per pass at 900°C or lower is 15% or more. Furthermore, it is preferable to include at least two rolling passes in which the reduction rate per pass at 900°C or lower is 15% or more. It is more preferable to include at least three rolling passes in which the reduction rate per pass at 900°C or lower is 15% or more, and even more preferable to include at least four rolling passes. Additionally, while there is no specific upper limit on the number of rolling passes in which the reduction rate per pass is 15% or more, it may be 50 passes or less, more preferable to 30 passes or less, and even more preferable to 15 passes or less. In addition, the reduction rate per pass at 900℃ or lower is not specifically limited to an upper limit, but it is preferable that it be 40% or less.

[0143] If the cumulative reduction rate at 900°C or lower is less than 50% or the rolling end temperature exceeds 850°C, the fineness of the austenite grains is insufficient, and low-temperature toughness deteriorates. Additionally, because the dislocation density within the grains decreases, the strength in the medium temperature range deteriorates. Furthermore, if the rolling end temperature is below the Ar3 temperature, a structure containing ferrite and bainite is formed, and the strength in the medium temperature range deteriorates. Therefore, the cumulative reduction rate at 900°C or lower is 50% or more, and the rolling end temperature is set from the Ar3 temperature to 850°C. The rolling end temperature is set to the Ar3 temperature or higher, preferably at the Ar3 temperature + 5°C or higher, more preferably at the Ar3 temperature + 10°C or higher, and even more preferably at the Ar3 temperature + 15°C or higher. In addition, the rolling end temperature is preferably 850°C or lower, 840°C or lower, more preferably 830°C or lower, and even more preferably 820°C or lower.

[0144] Accelerated cooling process

[0145] In the method for manufacturing a high-strength steel sheet according to the present invention, the accelerated cooling process is a process of accelerating cooling a hot-rolled steel sheet (hot-rolled sheet) obtained from a hot rolling process under conditions where the cooling start temperature is Ar3 or higher, the average cooling rate is 5℃ / s or higher, and the cooling stop temperature is 300 to 550℃. At this time, the average cooling rate refers to the average cooling rate from the cooling start temperature to the cooling stop temperature, and 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 cooling start to the cooling stop.

[0146] To suppress the formation of ferrite on the front and back surfaces of the steel sheet and to increase the bainite fraction, the cooling start temperature is set to be above the Ar3 temperature. Furthermore, it is preferable that the cooling start temperature be above the Ar3 temperature + 10°C, more preferable that it be above the Ar3 temperature + 15°C, even more preferable that it be above the Ar3 temperature + 20°C, and most preferable that it be above the Ar3 temperature + 25°C. To prevent low-temperature toughness deterioration caused by an increase in the average crystal grain size, the upper limit is preferably set to below the Ar3 temperature + 150°C, more preferable that it be below the Ar3 temperature + 140°C, and even more preferable that it be below the Ar3 temperature + 130°C.

[0147] The strength of high-strength steel sheets tends to increase with increasing average cooling rate during accelerated cooling. If the average cooling rate during accelerated cooling is less than 5°C / s, transformation begins at high temperatures; consequently, in addition to the formation of ferrite or pearlite in addition to bainite, the recovery of dislocations also proceeds during cooling. For this reason, if the average cooling rate is less than 5°C / s, sufficient strength cannot be obtained in the room temperature and medium temperature ranges. Furthermore, if the average cooling rate is less than 5°C / s, the effect of fine-graining the microstructure is reduced, the crystal grain size does not decrease, and low-temperature toughness deteriorates. Therefore, the average cooling rate during accelerated cooling is set to 5°C / s or higher. Additionally, it is preferable that the average cooling rate during accelerated cooling be 8°C / s or higher, more preferable that it be 10°C / s or higher, even more preferable that it be 15°C / s or higher, and most preferable that it be 20°C / s or higher. In addition, although the upper limit of the cooling rate is not specifically limited, in order to avoid an excessive increase in the martensite fraction, it is preferable that the average cooling rate be 80℃ / s or less, and more preferable that it be 50℃ / s or less.

[0148] The strength of steel sheets tends to increase as the cooling stop temperature of accelerated cooling decreases. If the cooling stop temperature of accelerated cooling exceeds 550°C, the growth of carbides is promoted, and the amount of dissolved carbon decreases; consequently, sufficient strength, particularly in the medium temperature range, cannot be obtained in the high-strength steel after cooling and in the state of steel pipes made of such high-strength steel. For this reason, the cooling stop temperature of accelerated cooling is set to 550°C or lower. Preferably, it is 540°C or lower, more preferably 530°C or lower, and even more preferably 520°C or lower. On the other hand, if the cooling stop temperature is below 300°C, the formation of low-temperature transformation products such as martensite, which has many mobile dislocations, becomes significant. As a result, the recovery of dislocations is promoted by prolonged aging in the medium temperature range, causing a significant decrease in strength. Therefore, the cooling stop temperature of accelerated cooling is set to 300°C or higher. Preferably 310°C or higher, more preferably 320°C or higher, even more preferably 330°C or higher, and most preferably 340°C or higher.

[0149] Method for manufacturing steel pipes

[0150] The method for manufacturing a steel pipe according to the present invention comprises a cold forming process and a welding process.

[0151] Cold forming process

[0152] The cold forming process is a process of cold forming the high-strength steel plate of the present invention into a tubular shape. When manufacturing a steel pipe for steam transport, the thickness of the steel plate is preferably 12 mm or more. In addition, it is preferably 30 mm or less. A more preferred range is the same as the plate thickness of the steel pipe described above. The method of cold forming the steel plate into a tubular shape is not particularly limited. Examples of forming methods include UOE forming, press bend forming, roll forming, etc.

[0153] Welding process

[0154] The welding process is a process of welding the butt joints formed by joining the ends of steel plates that have been shaped into a tube during the cold forming process. The welding method is not particularly limited, but it is preferable to use submerged arc welding or similar methods for the welded joint. Additionally, it is desirable to expand the steel pipe after welding, as this improves the roundness of the pipe cross-section. Heat treatment after manufacturing the steel pipe may be performed according to the desired characteristics and is not specifically prescribed.

[0155] Examples

[0156] Steel plates (thickness 15–25 mm) were cold-formed using steels A–AH having the compositional composition shown in Table 1 under the manufacturing conditions shown in Tables 2-1 and 2-2, followed by seam welding, and then 1.0% expansion was performed to manufacture steel pipes with an outer diameter of 610 mm × a pipe thickness of 15–25 mm. In addition, in the manufacturing conditions shown in Tables 2-1 and 2-2, "reduction rate" refers to the cumulative reduction rate at 900°C or lower, "FT" refers to the rolling end temperature, and "heat treatment" refers to long-term aging treatment. A 1.0% expansion rate means widening the inner diameter of the pipe by 1.0% using an expander in the pipe diameter direction from the inner side to the outer side of the steel pipe.

[0157] Using the methods described below, microstructure observation of steel plates, low-temperature toughness evaluation, tensile tests before and after aging, and tensile tests of steel pipes before and after aging were performed.

[0158] Microstructure observation of steel plates

[0159] A sample for observing the steel structure was taken from the center of the plate width of the steel plate manufactured as described above, and after mirror polishing the L cross-section of the steel plate (a cross-section parallel to the rolling direction and parallel to the normal direction of the rolling plane), the crystal orientation of a randomly selected 1 mm × 1 mm area at the center of the plate thickness was measured by electron backscatter diffraction (EBSD), and the crystal grain size of bainite was determined by image analysis using the area where the angle difference between adjacent pixels is 15° or more as the crystal grain boundary.

[0160] In addition, the average crystal grain size darea (㎛) is the area ai (㎛) occupied by each crystal grain. 2 It was calculated from the original diameter di (㎛) of each crystal grain and the following formula.

[0161] darea(㎛)=Σ(ai·di) / Σai

[0162] In addition, the top 20% of the crystal grain sizes of bainite represent the grain size of the 20th largest crystal grain when the equivalent diameters of each crystal grain are arranged in order of largest to smallest, and 20% of the total number of crystal grains are selected from the largest.

[0163] In addition, the microstructure was revealed by mirror polishing the L-section of the steel plate (a section parallel to the rolling direction and the normal direction of the rolling plane) and then performing Nital etching. Subsequently, using an optical microscope, a randomly selected 7.1 × 10⁻⁶ section at the center of the plate thickness was examined. －2 The steel structure was photographed over five fields of view in an area of ​​mm² (magnification: 400x), and the bainite fraction, ferrite fraction, and island-shaped martensite fraction within the photographed fields of view were used as area ratios, and the bainite grain size was measured using an image analysis device (Fiji) and an image analysis device (OTM Analysis manufactured by TSL). The average values ​​of the five fields of view were used as the bainite fraction, ferrite fraction, island-shaped martensite fraction, and bainite grain size.

[0164] In addition, the number of precipitates is determined using the method evaluated in the embodiment, as described above.

[0165] Low-temperature toughness evaluation of steel plates

[0166] In addition, the low-temperature toughness of the steel sheet was evaluated by conducting a DWTT in accordance with API 5L. The DWTT specimens were taken such that the thickness direction of the steel sheet coincided with the thickness direction of the specimen and the width direction of the steel sheet coincided with the length direction of the specimen. Furthermore, a ductile fracture rate of 85% or more and a fracture transition temperature of -40°C or lower at a test temperature of -40°C were evaluated as good.

[0167] Tensile test of steel plates before and after aging

[0168] As for the characteristics of the steel plate (before aging), a tensile test specimen was taken such that the direction perpendicular to the steel plate rolling direction and the direction perpendicular to the plate thickness direction was the length direction of the tensile test specimen, and a tensile test was performed at 350°C to determine the yield strength. In the tensile test, a round bar specimen with a diameter of 6 mm was used, and the crosshead speed was set to 0.15 mm / min. A yield strength of 555 MPa or higher in the tensile test at 350°C was evaluated as good.

[0169] As a characteristic of the long-term aging process (heat treatment at 400°C for 2335 hours), a tensile test specimen was taken such that the direction perpendicular to the steel plate rolling direction and perpendicular to the plate thickness direction was the length direction of the specimen, and heat treatment was performed at 400°C for 2335 hours in an N (nitrogen) atmosphere furnace, and a tensile test was conducted at 350°C to evaluate the yield strength. In the tensile test, a round bar specimen with a diameter of 6 mm was used, and a crosshead speed of 0.15 mm / min. A yield strength of 555 MPa or higher in the tensile test at 350°C was evaluated as good.

[0170] In addition, regarding the change in yield strength when maintained in the medium temperature range for a long time, the difference in yield strength before and after the aforementioned long-term aging, i.e., (yield strength before aging – yield strength after aging), was evaluated.

[0171] In addition, regarding the characteristics after long-term aging (heat treatment at 400°C for 2335 hours), for the steel pipe, a 6 mm diameter round bar test specimen was taken so that the circumferential direction of the pipe was the length direction of the test specimen, and heat treatment was carried out at 400°C for 2335 hours in an N (nitrogen) atmosphere in a furnace, and a tensile test was performed at 350°C, and the yield strength was evaluated as good when the yield strength at 350°C was 555 MPa or higher.

[0172] As described above, Tables 2-1 and 2-2 show the manufacturing conditions of the steel plate and the test results of the steel plate and steel pipe.

[0173] Examples of the present invention (No. 1 to 29), in which both the composition and the steel plate manufacturing conditions are within the scope of the present invention, have a yield strength of 555 MPa or more at 350°C for steel plates and steel pipes, and also have a yield strength of 555 MPa or more at 350°C after long-term aging of steel plates and steel pipes, and (yield strength before aging - yield strength after aging) is 50 MPa or less, and also have a ductile fracture rate of 85% or more and a fracture transition temperature of -40°C obtained by DWTT (test temperature: -40°C) in accordance with API 5L, and a fracture transition temperature of -40°C or less.

[0174] Meanwhile, comparative examples (No. 30 to 59), in which the composition or steel plate manufacturing conditions are outside the scope of the present invention, had at least one of the following: a yield strength at 350°C and / or a yield strength at 350°C before and after long-term aging of 555 MPa or more, a ductile fracture rate of 85% or more obtained by DWTT (test temperature: -40°C), and a fracture transition temperature of -40°C or less, lagged behind the present invention example.

[0175] [Table 1]

[0176]

[0177] [Table 2-1]

[0178]

[0179] [Table 2-2]

[0180]

Claims

Claim 1 The composition of the components is, in mass%, C: 0.04–0.09%, Si: 0.03–0.25%, Mn: 1.5–2.5%, P: 0.020% or less, S: 0.002% or less, Mo: 0.10–0.50%, Nb: 0.010–0.055%, Ti: 0.005–0.020%, Ca: 0.0040% or less, Al: 0.01–0.04%, N: 0.006% or less, the remainder being Fe and unavoidable impurities, X(%) represented by Equation (1) is 0.65% or more, the microstructure has bainite with an area ratio of 80% or more at the center of the plate thickness, the average crystal grain size of the bainite is 25㎛ or less, and among the top 20% of the crystal grains with large grain sizes among the bainite A high-strength steel sheet having a minimum crystal grain size of 60 μm or less, a ductile fracture rate of 85% or more obtained by DWTT at -40°C, and a yield strength of 555 MPa or more before and after aging under the condition of LMP = 15700, a Larson Miller parameter defined by the following equation (2). X(%) = 0.35Cr + 0.9Mo + 12.5Nb + 8V···(1) The element symbols in equation (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) Claim 2 A high-strength steel sheet according to claim 1, wherein the above component composition additionally contains, in mass%, one or two or more of Cr: 0.50% or less, V: 0.070% or less, Cu: 0.50% or less, and Ni: 0.50% or less. Claim 3 A high-strength steel sheet according to claim 1 or 2, additionally having ferrite with an area ratio of 10% or less and island-shaped martensite with an area ratio of 10% or less in the center of the sheet thickness, and having a total of 50,000 to 1,000,000 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, and composite precipitates with a diameter of 100 nm or less containing two or more elements among Ti, Nb, V, Mo, Cr, and Al in the center of the sheet thickness, and having a difference in yield strength before and after aging, calculated by subtracting the yield strength after aging from the yield strength before aging, of 50 MPa or less. Claim 4 A steel pipe using a high-strength steel plate as described in any one of paragraphs 1 to 3. Claim 5 A method for manufacturing a high-strength steel sheet as described in any one of claims 1 to 3, comprising: a heating process for heating a steel material to 1000 to 1200°C; a hot rolling process for hot rolling the steel material heated in the heating process, wherein the rolling includes at least one pass of rolling at 950°C or higher with a reduction rate of 10% or more per pass and at least one pass of rolling at 900°C or lower with a reduction rate of 15% or more per pass, wherein the cumulative reduction rate at 900°C or lower is 50% or more and the rolling end temperature is Ar3 or higher and 850°C or lower; and an accelerated cooling process for accelerating cooling the hot-rolled steel sheet obtained in the hot rolling process, wherein the cooling start temperature is Ar3 or higher, the average cooling rate is 5°C / s or higher, and the cooling stop temperature is 300 to 550°C. Claim 6 A method for manufacturing a steel pipe having a cold forming process for cold forming a high-strength steel plate into a tubular shape as described in any one of claims 1 to 3, and a welding process for welding a butting portion in which the ends of the steel plates formed into a tubular shape in the cold forming process are butted together.

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