JP4975304B2 - Method for producing high-strength steel sheet having high tensile strength of 760 MPa class or more excellent in hydrogen-induced crack resistance and ductile fracture characteristics, and method for producing high-strength steel pipe using the steel sheet - Google Patents

Description

  The present invention relates to a method for producing a high-strength steel plate having a tensile strength of 760 MPa or more and excellent in resistance to hydrogen-induced cracking, and a method for producing a high-strength steel pipe using the steel plate, suitable for natural gas / crude oil transportation line pipes, etc. About.

  In recent years, in pipelines for crude oil and natural gas, it has been required to reduce the outer diameter and weight of the line pipe for the purpose of increasing the internal pressure for the purpose of improving transportation efficiency and improving the local construction efficiency. Development of high-strength steel pipes exceeding 760 MPa to less than 900 MPa is underway (see, for example, Patent Documents 1 and 2).

  Generally, in pipelines, ductile cracks that occur in the base material of steel pipes can propagate over a long distance of 100 to several kilometers at a high speed of 100 m / s or more in the pipe axis direction. Required. The arrest resistance is the property that stops the propagation of cracks, the property that a brittle crack propagates through the base material and stops, that is, the property that the brittle fracture resistance and the property that the ductile crack propagates through the base material and stops. That is, it is classified as a ductile fracture characteristic.

  The brittle fracture resistance is evaluated by a drop weight tear test (hereinafter referred to as a DWTT test) and a temperature at which the ductile fracture surface ratio is 85% or more (referred to as a DWTT transition temperature). Brittle cracks often occur from welds, and steel pipes with excellent brittle fracture resistance evaluated by performing a DWTT test by introducing a weld bead at the center of the specimen and introducing a brittle crack have been proposed. (For example, refer to Patent Document 3).

  On the other hand, for evaluation of ductile fracture characteristics, a full crack burst test for determining whether or not a ductile crack generated by an explosion after mounting an explosive on the surface of a steel pipe stops is optimal. However, since the cost required for the full crack burst test is very high, conventionally, the Charpy absorbed energy or the absorbed energy required by the DWTT test (referred to as DWTT absorbed energy) that is relatively well consistent with the result of the full crack burst test. It is evaluated by. Thus, there has been a demand for the development of a high-strength steel plate and a high-strength steel pipe that are excellent in brittle fracture resistance and excellent in ductile fracture characteristics.

  On the other hand, in a high-strength steel plate or a high-strength steel pipe of X100 or higher, a defect called hydrogen induced cracking may occur. The hydrogen-induced crack is a crack that is generated parallel to the plate surface at the center of the plate thickness, and is caused by hydrogen. Since high-strength steel plates and high-strength steel pipes are highly sensitive to hydrogen, sometimes hydrogen-induced cracks exist in the steel plates and steel pipes. The presence of hydrogen-induced cracking has been a problem because it significantly deteriorates ductile fracture characteristics and crack propagation characteristics.

JP 09-041074 A Japanese Patent Application Laid-Open No. 09-041080 Japanese Patent Laid-Open No. 11-036042

  The present invention advantageously solves the above-mentioned problems, and has a high tensile strength in the circumferential direction of 760 MPa or higher (API standard X100 or higher) without hydrogen-induced cracking excellent in ductile fracture resistance and brittle fracture resistance. Provided is a method for producing a high-strength steel pipe having a tensile strength of 760 MPa or higher and excellent in hydrogen-induced crack resistance and ductile fracture characteristics, and a method for producing a high-strength steel pipe using the steel sheet, suitable for producing a high-strength steel pipe. is there.

  The present inventors have studied a high-strength steel sheet having a tensile strength in the circumferential direction of 760 MPa class or more and a method for preventing hydrogen-induced cracking of a steel pipe using the steel sheet, and a cooling method after stopping water cooling in steel sheet production As a result, an invention relating to a method for producing a high-strength steel plate free from hydrogen-induced cracking and a method for producing a high-strength steel pipe using the steel plate has been made. The gist of the present invention is as follows.

(1) By mass%, C: 0.01 to 0.5%, Si: 0.01 to 3.0%, Mn: 0.1 to 5.0%, P: 0.03% or less, S: When melting steel containing 0.03% or less, the balance being iron and inevitable impurities, the setting start temperature of the subsequent hot rolling is T1 (° C), the set finishing temperature is T3 (° C), When the set control cooling stop temperature after hot rolling is T4 (° C.), hydrogen as an impurity is H: {0.65+ (0.0007T4-0.03)} × 1.5 × in the molten steel. exp [-1411 {1 / (T1 + 273) -1 / (T3 + 273)}] The components are adjusted while being limited to ppm or less, the molten steel is cast, and further, hot rolling is performed at 1000 to 1250 ° C T1 (° C). Start, finish hot rolling at T3 (° C.) of 600 to 900 ° C., and stop controlled cooling during subsequent cooling The temperature T4 (° C.) is set to a temperature lower than the rolling end temperature of 50 ° C. or more, and controlled cooling is performed at a cooling rate of 0.5 to 20 ° C./s at the average cooling rate at the center of the steel sheet until the controlled cooling stop temperature T4 (° C.)該制allowed to cool from your cooling stop temperature T4 (° C.) to room temperature, characterized in that the amount of hydrogen of the steel sheet below 0.65 ppm, resistance to hydrogen induced cracking resistance and tensile excellent ductile fracture characteristic strength 760MPa grade The manufacturing method of the above high strength steel plate.

(2) By mass%, C: 0.01 to 0.5%, Si: 0.01 to 3.0%, Mn: 0.1 to 5.0%, P: 0.03% or less, S: When melting steel containing 0.03% or less, the balance being iron and inevitable impurities, the set heating temperature of the subsequent hot rolling is T2 (° C), the set finishing temperature is T3 (° C), When the set control cooling stop temperature after hot rolling is T4 (° C.), hydrogen as an impurity is H: {0.65+ (0.0007T4-0.03)} × 1.5 × in the molten steel. exp [-1411 {1 / (T2 + 273) -1 / (T3 + 273)}] The components are adjusted while being limited to ppm or less, and a steel slab obtained by casting the molten steel is heated to 1000 to 1250 ° C at a heating temperature T2 ( ℃), and after subsequent recrystallization zone rolling, at T3 (℃) of 600-900 ℃ During the subsequent cooling, the control cooling stop temperature T4 (° C.) is set to less than the rolling end temperature 50 ° C. or more, and the average cooling rate at the center of the steel sheet is 0.5 to the control cooling stop temperature T4 (° C.). controlled cooling at a cooling rate to be to 20 ° C. / s,該制allowed to cool from your cooling stop temperature T4 (° C.) to room temperature, characterized in that the amount of hydrogen of the steel sheet below 0.65 ppm, anti hydrogen induced A method for producing a high-strength steel sheet having a tensile strength of 760 MPa or more and excellent in crackability and ductile fracture characteristics.

(3) Instead of the steel components, in mass%, C: 0.02 to 0.10%, Si: 0.01 to 0.6%, Mn: 1.5 to 2.5%, P: 0 0.015% or less, S: 0.003% or less, Ni: 0.1 to 2.0%, Mo: 0.15 to 0.60%, Nb: 0.001 to 0.10%, Ti: 0.0. 005 to 0.030%, Al: 0.06% or less, N: 0.0001 to 0.006%, characterized by containing hydrogen-resistant cracking according to (1) or (2) above For producing high-strength steel sheets having a tensile strength of 760 MPa or more and excellent in ductility and ductile fracture characteristics.

(4) Further, by mass%, B: 0.0001 to 0.005%, V: 0.001 to 0.10%, Cu: 0.01 to 1.0%, Cr: 0.01 to 0.00. 8%, Zr: 0.0001 to 0.005%, Ta: 0.0001 to 0.005%, Ca: 0.0001 to 0.01%, REM: 0.0001 to 0.01%, Mg: 0 It is excellent in hydrogen-induced crack resistance and ductile fracture characteristics according to any one of the above (1) to (3), characterized by containing .0001 to 0.006% of one or more. A method for producing a high-strength steel sheet having a tensile strength of 760 MPa or higher.

(5) The reheating temperature of the steel slab is 1000 to 1250 ° C., and is excellent in hydrogen-induced crack resistance and ductile fracture characteristics according to any one of (2) to (4) above. A method for producing a high-strength steel sheet having a tensile strength of 760 MPa or higher.

(6) The hydrogen-induced crack resistance and ductile fracture characteristics according to any one of (1) to (5) above, wherein the steel sheets are stacked and allowed to cool after the controlled cooling is stopped. A method for producing a high-strength steel sheet having an excellent tensile strength of 760 MPa or higher.

(7) The laminated steel sheet is covered with a heat insulating cover and allowed to cool, and the tensile strength of 760 MPa class or more excellent in hydrogen-induced crack resistance and ductile fracture characteristics according to (6) above Manufacturing method of steel sheet.

(8) It is excellent in hydrogen-induced crack resistance and ductile fracture characteristics, characterized by being piped using a high-strength steel sheet produced by the method described in any one of (1) to (7) above. A method for producing a high strength welded steel pipe having a tensile strength of 760 MPa or higher.

(9) In the pipe making process, the high-strength steel sheet is formed into a tubular shape in the UO process, the ends are submerged arc welded using a welding wire and a fired flux or a molten flux, and then the pipe expansion is performed. The method for producing a high-strength welded steel pipe having a tensile strength of 760 MPa or more and excellent in hydrogen-induced crack resistance and ductile fracture characteristics as described in (8) above.
(10) By mass%, C: 0.01 to 0.12%, Si: 0.3% or less, Mn: 1.2 to 2.4%, Ni: 4.0 to 8.5%, Cr + Mo + V: Submerged arc welding is performed using a welding wire containing 3.0 to 5.0%, Ti: 0.005 to 0.15%, Al: 0.02% or less, with the balance being iron and inevitable impurities. The method for producing a high-strength welded steel pipe having a tensile strength of 760 MPa or more and excellent in hydrogen-induced cracking resistance and ductile fracture characteristics according to (9) above

(11) The hydrogen resistance according to (9) or (10) above, wherein the heat input per 1 mm thickness of the submerged arc welding is 0.13 to 0.25 kJ / mm ^2. A method for producing a high-strength welded steel pipe having a tensile strength of 760 MPa or more and excellent in induced cracking and ductile fracture characteristics.

  According to the present invention, a method for producing a high-strength steel sheet having a tensile strength of 760 MPa or higher (API standard X100 or higher) excellent in hydrogen-induced cracking resistance and ductile fracture characteristics, suitable for natural gas / crude oil transportation line pipes, and the like Since the manufacturing method of the high strength steel pipe using a steel plate can be provided, the industrial effect is immeasurable.

  Hereinafter, the present invention will be described in detail.

  The present inventors detail the relationship between the water cooling stop temperature of the steel sheet after hot rolling and the amount of hydrogen released when it is allowed to cool from the water cooling stop temperature to room temperature (hereinafter referred to as the amount of hydrogen released during cooling). Derived. FIG. 1 shows the relationship between the water cooling stop temperature and the amount of hydrogen released during cooling. The amount of hydrogen remaining in the steel at room temperature (hereinafter referred to as the residual hydrogen amount) is finally released from the amount of hydrogen in the steel immediately after hot rolling (hereinafter referred to as the hydrogen amount immediately after rolling). It can be considered as an amount obtained by subtracting the amount of hydrogen released during cold.

  Here, a method for deriving the amount of hydrogen immediately after rolling, the amount of hydrogen released during cooling, and the amount of residual hydrogen will be described. First, the hydrogen amount immediately after rolling is immediately cooled with water after rolling the steel plate, and the steel plate is inserted into liquid nitrogen. Thereafter, a 10 mm square × 40 mm long square test piece is produced while cooling with dry ice. The temperature of the test piece is raised, and the “amount of hydrogen released during the temperature rise” is analyzed by a gas chromatography method (hereinafter referred to as a temperature rise desorption method). The amount of hydrogen analyzed here is defined as “the amount of hydrogen immediately after rolling”. Next, each steel plate that has been allowed to cool from the water cooling stop temperature to room temperature for each water cooling stop temperature (control cooling stop temperature) condition is inserted into liquid nitrogen with respect to the amount of hydrogen released during cooling. Thereafter, a 10 mm square × 40 mm long square test piece is produced while cooling with dry ice. The test piece for each water cooling stop temperature condition is analyzed for hydrogen content by the temperature programmed desorption method. The amount of hydrogen thus obtained is defined as “amount of released hydrogen during cooling” for each water cooling stop temperature (control cooling stop temperature) condition. Finally, the “residual hydrogen amount” for each water cooling stop temperature (control cooling stop temperature) condition is derived by subtracting each “released hydrogen amount during cooling” from the “hydrogen amount immediately after rolling”.

  Moreover, the measuring method of the amount of hydrogen at the time of melting of steel is measured by a combustion method. That is, a sample for analysis is taken from molten steel, processed in dry ice to a size of 3 mm × 3 mm × 10 mm, and immediately heated to 2000 ° C. At this time, the amount of hydrogen that has come out as gas is measured by mass spectrometry, and this measured value is referred to as “amount of hydrogen in molten steel”.

  As shown in FIG. 1, when the water cooling stop temperature (control cooling stop temperature) increases, the amount of hydrogen released during cooling increases, and conversely, the amount of residual hydrogen decreases. At this time, the relationship between the amount of hydrogen released during cooling (E) and the controlled cooling stop temperature (T4) is expressed by equation (1).

E = 0.007T4-0.03 (1)
Here, the unit of E is ppm, and the unit of T is ° C.

  Next, the crack limit hydrogen amount was derived using various steel plates for X100 line pipe that were accelerated to 0.5 to 20 ° C./s to room temperature. FIG. 2 shows the relationship between the residual hydrogen amount and the hydrogen-induced crack area ratio. The hydrogen-induced crack area ratio is represented by the area ratio of hydrogen-induced cracks existing in an area of 300 mm width × 300 mm length of the steel sheet. This hydrogen-induced crack area ratio is represented by image analysis as the area of a defect detected from the plate surface with an ultrasonic flaw detector. However, the detection limit length of the ultrasonic flaw detector is 1 mm. From FIG. 2, it was found that the crack limit hydrogen content of the steel plate for X100 line pipe was 0.65 ppm.

  Since the crack limit hydrogen amount was 0.65 ppm, the controlled cooling stop temperature (T4) was controlled at a temperature of C4 when taking into account the amount of hydrogen released during cooling as it was released during the cooling after rolling shown in Equation (1). The hydrogen cracking limit hydrogen amount (C1) of the steel sheet immediately after the cooling is finished is expressed by the equation (2).

C1 = 0.65 + (0.0007T4-0.03) (2)
Here, the unit of C1 is ppm, and the unit of T4 is ° C.

  For example, according to equation (2), the higher the control cooling stop temperature, the greater the crack limit hydrogen amount of the steel sheet immediately after the control cooling stop. For example, if the controlled cooling stop temperature is 500 ° C., the hydrogen cracking limit hydrogen amount of the steel plate at 500 ° C. is 1.0 ppm.

  Next, a method for deriving the amount of hydrogen in steel in the heating temperature or the hot rolling start temperature range will be described. From the temperature range of the heating temperature T2 (° C.) to the hot rolling start temperature T1 (° C.) to the finish rolling end T3 (° C.), it is considered that the amount of hydrogen present in the steel follows an equilibrium relationship with the rolling temperature. . The amount of hydrogen in the steel at each temperature T (° C.) from the temperature range of the heating temperature T 2 (° C.) to the hot rolling start temperature T 1 (° C.) to the hot rolling finish (end) temperature T 3 (° C.) is (3 ) Expression.

log (H) = − 1411 {1 / (T + 273)} − 0.468 (3)
For example, if the heating temperature is T2 (° C.) and the finish rolling temperature is T3 (° C.), the ratio of the amount of hydrogen present in the steel at T2 (° C.) to the amount of hydrogen present in T3 (° C.) is ( 4) It is expressed as shown below.

H _T2 / H _T3 = exp [-1411 {1 / (T2 + 273) −1 / (T3 + 273)}]
(4)
When the formulas (2) and (4) are combined, the cracking limit hydrogen amount (C2) in the heating temperature range is expressed as the formula (5).

C2 = {0.65+ (0.0007T4-0.03)} × 1.5 × exp [−1411 {1 / (T2 + 273) −1 / (T3 + 273)}] (5)
Here, the correction coefficient 1.5 is a correction value that takes into account the amount by which the amount of hydrogen in the steel decreases during the γ → α transformation. FIG. 3 shows the relationship between the cracking limit hydrogen amount during heating and the heating temperature (T2) when the hot rolling finish (end) temperature (T3) is 800 ° C. and the controlled cooling stop temperature (T4) is 500 ° C. As shown in FIG. 3, for example, the crack limit hydrogen amount at a heating temperature of 1200 ° C. is 2.2 ppm, and at a heating temperature of 1000 ° C., it is 1.9 ppm.

  Here, when the steel is melted, cast, and then directly rolled, the crack limit hydrogen amount (C3) is the rolling start temperature T1 (° C.) instead of the reheating temperature T2 (° C.) in the equation (5). Instead, it can be rewritten as in equation (6).

C3 = {0.65+ (0.0007T4-0.03)} × 1.5 × exp [−1411 {1 / (T1 + 273) −1 / (T3 + 273)}] (6)
Therefore, at the time of melting the steel, as described above, it is sufficient to control the amount of hydrogen so that it is not more than the crack limit hydrogen amount in the heating temperature in the subsequent hot rolling or the hot rolling start temperature range. become.

  Next, a method for increasing the amount of hydrogen released after the controlled cooling is stopped after the hot rolling is completed will be described. As a result of examining various methods for releasing a large amount of hydrogen during cooling after rolling more than the value obtained by the formula (1), when steel plates are stacked, the cooling time to room temperature becomes longer, and rolling It was found that the amount of hydrogen released during subsequent cooling increases. Furthermore, after stacking the steel plates, it is found that if these steel plates are covered with a heat insulating cover, the time for cooling to room temperature becomes longer, and in this case also, the amount of hydrogen released during cooling is increased.

  Next, the reason for limiting the chemical component of the base material will be described.

  C is an extremely effective element for improving the strength of steel, and in order to obtain the strength of the present invention, it is necessary to contain 0.01% or more. Furthermore, it is preferable to contain 0.02% or more of C. This is because when the content is 0.02% or more, it is easy to ensure the strength. However, if the C content is more than 0.5%, the low temperature toughness of the base metal and the weld heat-affected zone (referred to as HAZ) may be slightly deteriorated, which may impair the on-site weldability. 5% or less. Furthermore, it is preferable that the upper limit of the C content is 0.1% or less. This is because if it exceeds 0.1%, the weldability tends to deteriorate.

  Si is an element effective for deoxidation, and in order to obtain the effect, it is necessary to contain 0.01% or more. On the other hand, if the content is more than 3.0%, the low temperature toughness of the HAZ is slightly deteriorated and the on-site weldability may be impaired, so the upper limit of the Si content is set to 3.0%. Furthermore, it is preferable to set it as 0.6% or less. This is because if it exceeds 0.6%, the HAZ toughness tends to deteriorate.

  Mn is an element effective for improving the balance between the strength and low temperature toughness of steel, and in order to obtain the effect, the lower limit of the Mn content needs to be 0.1% or more. However, if Mn is contained in excess of 5.0%, the hardenability of the steel is increased and the low temperature toughness of the HAZ is deteriorated, and the field weldability may be impaired. And Furthermore, it is preferable that the upper limit of the Mn content is 2.5% or less. This is because if it exceeds 2.5%, the low temperature toughness and the HAZ toughness tend to deteriorate.

  P and S are impurity elements, and in order to further improve the low temperature toughness of the base material and the HAZ, the upper limits of the P content and the S content are 0.03% or less and 0.03% or less, respectively. There is a need to. Furthermore, it is desirable to make it 0.015% or less and 0.003% or less, respectively. This is because when it exceeds 0.015% and 0.003%, low temperature toughness, HAZ toughness and hydrogen-induced cracking are likely to occur. Although the lower limit of the P content and the S content is preferably as low as possible, it is not specified, but usually contains 0.001% or more and 0.0001% or more, respectively, in terms of steelmaking capacity.

  Ni is an element that improves low-temperature toughness and strength, and in order to obtain the effect, the lower limit of the Ni content is preferably set to 0.1% or more. On the other hand, if the Ni content exceeds 2.0%, weldability may be impaired, so the upper limit of the Ni content is preferably 2.0%.

  Mo is an element that improves the hardenability of steel and forms carbonitride to improve strength. To obtain the effect, Mo content is preferably 0.15% or more. On the other hand, if the Mo content exceeds 0.60%, the strength becomes too high and the low temperature toughness of the HAZ may be impaired, so the upper limit of the Mo content is preferably 0.60%.

  Nb is an element that forms carbides and nitrides and improves the strength of steel. To obtain this effect, the Nb content is preferably 0.001% or more. On the other hand, if the Nb content is more than 0.10%, the low temperature toughness of the base material and the HAZ may be impaired, so the upper limit of the Nb content is preferably 0.10%.

  Ti is an element that is effective for deoxidation and contributes to the refinement of the crystal grain size by forming a nitride. To obtain the effect, it is preferable to add 0.005% or more. On the other hand, if the Ti content is more than 0.030%, coarse carbides may be produced and the low temperature toughness may be deteriorated. Therefore, the upper limit of the Ti content is preferably 0.030% or less.

  Al is an effective element as a deoxidizer, but if the Al content exceeds 0.06%, Al-based non-metallic inclusions may increase and inhibit the cleanliness of the steel. The upper limit of 0.06% or less. In addition, since deoxidation is possible also with Ti and / or Si, it is not necessary to contain Al, Therefore A minimum in particular is not prescribed | regulated.

  N forms nitrides with Ti, Al, etc., and prevents the austenite grains in the weld heat affected zone from becoming coarse. This effect becomes prominent with addition of 0.0001% or more, but addition exceeding 0.006% may lead to a decrease in toughness. Therefore, it is preferable that the addition amount of N is in the range of 0.0001 to 0.006%.

  Furthermore, in the present invention, one or more elements of B, V, Cu, Cr, Zr, Ta, Ca, REM, and Mg can be added as elements for improving strength and toughness.

  B is an element that enhances hardenability and improves the toughness of the weld heat affected zone. This effect becomes prominent with addition of 0.0001% or more, but addition exceeding 0.005% may lead to a decrease in toughness. Therefore, it is preferable that the addition amount of B is in the range of 0.0001 to 0.005%.

  V is an element that forms carbides and nitrides in the same manner as Nb and improves the strength of the steel. However, in order to obtain a remarkable effect, V is preferably added in an amount of 0.001% or more. On the other hand, if V is added in excess of 0.10%, the toughness may be lowered, so the upper limit is preferably made 0.10% or less.

  Cu is an element that increases the strength, and is preferably added in an amount of 0.01% or more. On the other hand, if more than 1.0% is added, cracking is likely to occur during heating of the steel slab or during welding, so the upper limit is preferably made 1.0% or less.

  Cr is an element that improves the strength of steel by precipitation strengthening, and the addition of 0.01% or more is effective. On the other hand, if added in a larger amount than 0.8%, the hardenability of the steel is increased and the toughness may be lowered, so the upper limit is preferably made 0.8% or less.

  Zr and Ta are elements that form carbides and nitrides as in Nb and improve the strength of the steel, and are each preferably added in an amount of 0.0001% or more. On the other hand, when Zr and Ta are added in excess of 0.005%, toughness may be reduced. Therefore, it is preferable that the upper limit of the addition amount of Zr and Ta is 0.005% or less, respectively.

  Ca and REM suppress the production | generation of the extended | stretched MnS by producing | generating a sulfide, and improve the characteristic of the plate | board thickness direction of steel materials, especially lamella tear resistance. In order to obtain this effect, it is preferable to add 0.0001% or more of Ca and REM, respectively. On the other hand, when Ca and REM are added in excess of 0.01%, Ca and REM oxides increase. Therefore, it is preferable that the upper limit of the addition amount of Ca and REM is 0.01% or less, respectively.

  Mg is an element that generates fine Mg-containing oxides or sulfides such as MgO and MgS, suppresses coarsening of austenite grains, and improves HAZ toughness. In order to acquire this effect, it is preferable to add 0.0001% or more of Mg. On the other hand, if Mg is added in excess of 0.006%, the Mg-containing oxide and sulfide are coarsened, so the upper limit is preferably made 0.006% or less.

  Next, the manufacturing method of a steel plate is demonstrated. The steel containing the components shown above is melted in the steel making process, continuously cast, then heated, and hot rolled. Here, after continuous casting, the steel slab may be directly rolled without being cooled and reheated. This is because even if direct rolling is performed, the hydrogen release amount is not affected.

  The setting start temperature (T1) of hot rolling is 1000 to 1250 ° C. If the temperature is lower than 1000 ° C., sufficient recrystallization rolling may not be performed, and if it exceeds 1250 ° C., coarse grains may remain during rolling. The set heating temperature (T2) is set to 1000 to 1250 ° C. similarly to the set start temperature (T1) of hot rolling. If the temperature is lower than 1000 ° C., sufficient recrystallization rolling may not be performed, and if it exceeds 1250 ° C., coarse grains may remain during rolling.

  On the other hand, the set finishing (end) temperature (T3) of hot rolling is set to 600 to 900 ° C. If the temperature is lower than 600 ° C., the strength is high and the rolling load is too large, so the shape after rolling cannot be controlled. Further, if the temperature exceeds 900 ° C., the γ grains do not become fine, so the temperature is set to 900 ° C. or less.

  Further, when cooling after the end of hot rolling, the control cooling stop temperature T4 (° C.) is set to 50 ° C. or more below the rolling end temperature, and the average cooling rate at the center of the steel sheet is set to 0. 0 to the control cooling stop temperature T4 (° C.). Control cooling is performed at a cooling rate of 5 to 20 ° C./s. If the cooling rate of the controlled cooling is less than 0.5 ° C./s, the strength cannot easily satisfy 760 MPa, and if it exceeds 20 ° C./s, the strength of the steel sheet is too high and the steel pipe is formed into a steel pipe. Since it may not be possible, it is set to 0.5 to 20 ° C./s. If the controlled cooling stop temperature T4 (° C.) is less than 50 ° C., the amount of hydrogen released after the controlled cooling stop is eliminated, so the lower limit is set to 50 ° C. In addition, the controlled cooling is preferably water cooling that is easy to control. For example, in a steel sheet having a thickness of 20 mm or less, a cooling rate of 0.5 to 20 ° C./s may be ensured even by cooling other than water cooling. Needless to say, it is not limited to water cooling.

  Next, the preferable range of the seam weld metal component when the steel plate is a steel pipe will be described.

  C is extremely effective for improving the strength of the steel, and in order to obtain the target strength in the martensite structure, the C content is preferably 0.04% or more. On the other hand, if the C content exceeds 0.14%, welding low temperature cracking is likely to occur, and the on-site welded part and seam welding intersect, leading to an increase in the HAZ maximum hardness of the so-called T-cross part. The upper limit is preferably 0.14% or less. Furthermore, the upper limit of preferable C content is 0.10% or less. This is because if it exceeds 0.1%, the toughness tends to deteriorate.

  In order to prevent the occurrence of blow holes, Si is preferably contained in an amount of 0.05% or more. On the other hand, when the Si content is more than 0.4%, the low temperature toughness may be deteriorated, and particularly when performing inner and outer surface welding or multilayer welding, the low temperature toughness of the reheated portion may be deteriorated. The upper limit is preferably 0.4% or less.

  Mn is an element that improves the balance between strength and low-temperature toughness and forms inclusions that form nuclei for intragranular bainite. In order to obtain this effect, the Mn content is preferably 1.2% or more. On the other hand, if the Mn content is more than 2.2%, segregation is promoted and the low temperature toughness may be deteriorated, making it difficult to produce a welding material. Therefore, the upper limit of the Mn content is 2.2% or less. It is preferable that

  P and S are inevitable impurities, and in order to suppress deterioration of low temperature toughness and reduce low temperature cracking susceptibility, the smaller the amount, the smaller the content of P and S, 0.01% or less, respectively. It is preferable to set it to 01% or less.

  Ni is an element that enhances hardenability and improves strength and improves low-temperature toughness. In order to obtain this effect, it is preferable to contain 1.3% or more of Ni. On the other hand, if the Ni content is more than 3.2%, hot cracking may occur, so the upper limit of the Ni content is preferably 3.2% or less.

  Cr, Mo, and V are all elements that increase the hardenability and improve the strength. To obtain the effect, Cr + Mo + V is preferably set to 1.0% or more. On the other hand, if Cr + Mo + V is added in a larger amount than 2.5%, low temperature cracking may occur, so the upper limit of the Cr + Mo + V content is preferably 2.5% or less.

  Ti is an element that forms a nitride, oxide, or the like of Ti that serves as a nucleus for formation of intragranular bainite, and is preferably contained in an amount of 0.003% or more. On the other hand, when the Ti content is more than 0.05%, a large amount of Ti carbide is generated and the low temperature toughness may be deteriorated. Therefore, the upper limit of the Ti content is preferably 0.05%.

  Since Al sometimes inhibits the formation of Ti oxides that form nuclei of intragranular bainite, it is preferable that the Al content is low. The upper limit with preferable Al content is 0.02% or less, More preferably, 0.015% or less is good. This is because if it exceeds 0.02% and 0.015%, the toughness tends to deteriorate and low temperature cracking tends to occur.

  B is an element that enhances the hardenability and improves the low temperature toughness of the weld metal, and preferably contains 0.0003% or more, but if the B content is more than 0.005%, the low temperature toughness is reduced. Since it may deteriorate, the B content is preferably 0.005% or less.

  O is an element that lowers the hardenability and degrades the low temperature toughness of the weld metal. When the amount of O exceeds 0.06%, the low temperature toughness is remarkably deteriorated. On the other hand, if the amount of O is low, cold cracking is likely to occur, and at the same time, the on-site weldability deteriorates, so 0.010% or more is preferable.

  In addition, the weld metal may contain elements such as Zr, Nb, and Mg that are added to improve the refining and solidification during welding.

  The structure of the weld metal is mainly composed of bainite / martensite and intragranular bainite, and the balance is ferrite and / or retained austenite. In order to set the tensile strength to 760 MPa or more, it is preferable to set the area ratio of bainite / martensite to 50% or more.

  Further, in order to improve the low temperature toughness of the weld metal, it is preferable that the area ratio of intragranular bainite is large, and it is preferable to set it to 10% or more. Bainitic martensite and intragranular bainite can be distinguished by structural observation with an optical microscope or scanning electron microscope, and the area ratio of bainite martensite and intragranular bainite was photographed with an optical microscope or scanning electron microscope. It can be measured by image analysis using tissue photographs.

  Furthermore, the steel plate is press-formed into a cylindrical shape, and the ends are submerged arc welded to form a steel pipe.

  Submerged arc welding is a welding with a large dilution of the base metal, and in order to obtain a desired characteristic, that is, a weld metal composition, it is necessary to select a welding material in consideration of the dilution of the base metal. Hereinafter, although the reason for limiting the chemical composition of the welding wire will be described, it is basically a production method capable of realizing a high-strength line pipe having a tensile strength of 760 MPa or higher.

  In order to obtain the C content in the range required for the weld metal, C is set to 0.01 to 0.12% in consideration of dilution with the base material component and mixing of C from the atmosphere.

  Si, Mn, Ni, Cr + Mo + V are each 0.3% or less in consideration of dilution by the base material component in order to obtain the content of Si, Mn, Ni, Cr + Mo + V in the preferred weld metal component range. 2 to 2.4%, 4.0 to 8.5%, and 3.0 to 5.0%.

  Ti is an element that forms a nitride, oxide, or the like of Ti that forms nuclei for intragranular bainite, and it is preferable to contain 0.005% or more. On the other hand, if the Ti content is more than 0.15%, a large amount of Ti carbide is generated and the low temperature toughness may be deteriorated. Therefore, the upper limit of the Ti content is preferably set to 0.15%.

  Since Al sometimes inhibits the formation of Ti oxides that form nuclei of intragranular bainite, it is preferable that the Al content is low. The upper limit with preferable Al content is 0.02% or less.

  In addition, it is desirable that impurities of P and S are as small as possible, and B can be added to ensure strength. Moreover, Zr, Nb, Mg, etc. may be used for the purpose of deoxidation, and at least one or more of these elements may be added.

  In addition, welding can be performed with a plurality of electrodes as well as a single electrode. In the case of welding with a plurality of electrodes, it is possible to combine various wires, and it is not necessary for each wire to be in the above component range, and it is sufficient that the average composition from each wire component and consumption is in the above component range.

  Flux used for submerged arc welding can be broadly classified into fired flux and molten flux. Firing-type fluxes have the advantage that alloy materials can be added and the amount of diffusible hydrogen is low, but they have the disadvantage of being easily pulverized and difficult to use repeatedly. On the other hand, the melt-type flux is in the form of glass powder, has the advantages of high grain strength and is difficult to absorb moisture, and has the disadvantage that diffusible hydrogen is slightly high. When producing the high-strength steel pipe of the present invention, welding cold cracking is likely to occur. From this point, a firing mold is desirable, but a molten mold that can be recovered and used repeatedly is suitable for mass production and has a low cost. There are advantages. The cost is high in the baking mold, and the necessity of strict quality control is a problem in the melting mold, but it is within a range that can be handled industrially, and either can be used essentially.

  Next, welding conditions will be described below.

  The initial tack welding performed may be any of MAG arc welding, MIG arc welding, and TIG arc welding. Usually, MAG arc welding. Next, the inner and outer surfaces are preferably submerged arc welding, but may be TIG arc welding, MIG arc welding, or MAG arc welding. The inner and outer surfaces may be welded one by one, but a plurality of passes may be performed.

  When submerged arc welding is performed on the inner and outer surfaces, if the welding speed is less than 1 m / min, seam welding of the line pipe is inefficient, and if it exceeds 3 m / min, the bead shape may become unstable. Therefore, the welding speed of submerged arc welding is preferably in the range of 1 to 3 m / min.

In addition, when the welding part of tack welding and inner and outer surface welding overlaps, the one where welding heat input is as low as possible is preferable. In addition, although the welding heat input varies depending on the plate thickness, if the heat input is too small, the penetration becomes insufficient, the number of times of welding increases, and the working efficiency deteriorates. On the other hand, if the welding heat input is too large, the heat-affected zone is greatly softened and the toughness of the weld zone is also reduced. Therefore, the specific heat of the inner and outer surfaces per 1 mm of plate thickness is preferably 0.13 to 0.25 kJ / mm ² . For example, the welding heat input of the inner and outer surfaces with a plate thickness of 15 mm is 1.9 to 3.8 kJ / mm.

  After seam welding, roundness is improved by pipe expansion. In order to make a perfect circle, it is necessary to deform to the plastic region. In the case of the high-strength steel pipe of the present invention, it is preferable that the pipe expansion ratio expressed as a percentage obtained by dividing the difference between the circumference after pipe expansion and the circumference before pipe expansion by the circumference before pipe expansion is 0.5% or more. On the other hand, if the expansion ratio exceeds 2.0%, the toughness may deteriorate due to plastic deformation of both the base material and the welded portion. Therefore, the tube expansion rate is preferably in the range of 0.5 to 2.0%.

  Steels composed of the chemical components shown in Table 1 were melted and cast into steel pieces. These steel slabs were hot rolled at various temperatures T1, and after the hot rolling was completed at various temperatures T3, the average cooling rate was 5 ° C / s, and various controlled cooling stop temperatures of 550 ° C or lower. Water-cooled to T4. With these 20 mm thick steel plates, the amount of hydrogen in the steel and the presence or absence of hydrogen-induced cracking were investigated. As shown in Table 2, when the amount of hydrogen in the steel sheet was less than 1.0 ppm, hydrogen-induced cracking did not occur, and the precrack DWTT absorbed energy was 3000 J or more, indicating good ductile fracture characteristics. Furthermore, the DWTT ductile fracture area ratio was 85% or more, and good arrestability was exhibited.

  Steel made of the chemical components shown in Table 3 was melted and cast into steel pieces having a thickness of 240 mm. These steel slabs are heated to various heating temperatures T2 and hot-rolled at 900 ° C. or higher to a thickness of 59 to 80 mm in a recrystallization temperature region, and then hot-rolled in an unrecrystallized region to a thickness of 14 to 25 mm. It carried out from 0 degreeC to hot rolling completion temperature T3 of various temperatures. In addition, 900 degreeC or more is a recrystallization temperature range, and 880 degrees C or less is a non-recrystallization temperature range. After hot rolling, it was water cooled to various controlled cooling stop temperatures T4 of 1 ° C./s or more and 10 ° C./s or less and 450 ° C. or less at an average cooling rate. The obtained steel sheet was press-formed into a cylindrical shape and tack welded. Then, the inner and outer surfaces were submerged arc welded with a welding heat input of 2.5 to 3.5 kJ / mm, expanded, and 36 inches ( 913 mm diameter) and 16 mm thick steel pipe. Tables 6 and 7 (continued in Table 6) show manufacturing conditions, characteristics of the base material, test results, and the like at this time.

  A hydrogen content analysis test piece of 10 mm × 10 mm × 40 mm was taken from a base material of 1/2 t part of the obtained steel pipe, and hydrogen analysis in the steel was performed by a temperature programmed desorption method. The presence or absence of hydrogen-induced cracking in the steel pipe was measured using an ultrasonic flaw detector.

  Implementation No. 11 to 21 show examples of the present invention. As is clear from Tables 6 and 7 (continued in Table 6), these steel sheets had a hydrogen content in steel of 0.65 ppm or less, and no hydrogen-induced cracking occurred. As a result, the precrack DWTT absorption energy was 3000 J or more, indicating good ductile fracture characteristics. Furthermore, the DWTT ductile fracture area ratio was 85% or more, and good arrestability was exhibited. On the other hand, the implementation No. Reference numerals 22 to 30 show comparative examples deviating from the method of the present invention. That is, the implementation No. In Nos. 22 to 30, the amount of hydrogen in the steel sheet exceeds 0.65 ppm, and hydrogen-induced cracking occurs. As a result, the precrack DWTT energy was less than 3000 J, and the fracture surface ratio was also less than 85%.

  Tables 4 and 5 (continued in Table 4) show the components of the weld metal and the weld wire of the present invention and the comparative example for reference.

It is a figure which shows the relationship between the water cooling stop temperature after rolling (control cooling stop temperature) in the manufacturing method of the high strength steel plate by this invention, and the amount of hydrogen discharge at the time of cooling after that. It is a figure which shows the relationship between the residual hydrogen amount and the hydrogen induced crack area ratio in the high strength steel plate by this invention. It is a figure which shows the relationship between the crack limit hydrogen amount in a heating temperature range, and heating temperature in the high strength steel plate by this invention.

Claims (11)

% By mass
C: 0.01-0.5%
Si: 0.01-3.0%,
Mn: 0.1 to 5.0%,
P: 0.03% or less,
S: When melting steel containing 0.03% or less, the balance being iron and inevitable impurities, the setting start temperature of hot rolling in the subsequent process is T1 (° C.), and the setting finishing temperature is T3 (° C. ), When setting control cooling stop temperature after hot rolling is T4 (° C.), hydrogen as an impurity in molten steel,
H: {0.65+ (0.0007T4-0.03)} × 1.5 × exp [-1411 {1 / (T1 + 273) −1 / (T3 + 273)}] ppm The molten steel is cast, and further hot rolling is started at 1000 to 1250 ° C. T1 (° C.), hot rolling is finished at 600 to 900 ° C. T3 (° C.), and the subsequent cooling is controlled cooling stop temperature. T4 (° C.) is set to a temperature lower than the rolling end temperature of 50 ° C. or higher, and controlled cooling is performed at a cooling rate of 0.5 to 20 ° C./s at an average cooling rate at the center of the steel sheet to the controlled cooling stop temperature T4 (° C.). allowed to cool from the control cooling stop temperature T4 (° C.) to room temperature, characterized in that the amount of hydrogen of the steel sheet below 0.65 ppm, resistance to hydrogen induced cracking resistance and ductile fracture characteristics in excellent tensile strength 760MPa grade or higher Manufacturing method of high strength steel sheet. % By mass
C: 0.01-0.5%
Si: 0.01-3.0%,
Mn: 0.1 to 5.0%,
P: 0.03% or less,
S: When melting steel containing 0.03% or less, the balance being iron and inevitable impurities, the set heating temperature of the hot rolling in the subsequent step is T2 (° C.), and the set finishing temperature is T 3 (° C. ), When setting control cooling stop temperature after hot rolling is T4 (° C.), hydrogen as an impurity in molten steel,
H: {0.65+ (0.0007T4-0.03)} × 1.5 × exp [-1411 {1 / (T2 + 273) −1 / (T3 + 273)}] ppm The steel slab obtained by casting the molten steel is reheated at a heating temperature T2 (° C.) of 1000 to 1250 ° C., and after subsequent recrystallization zone rolling, at T3 (° C.) of 600 to 900 ° C. When the hot rolling is finished and the subsequent cooling is performed, the controlled cooling stop temperature T4 (° C.) is set to 50 ° C. or less below the rolling finish temperature, and the average cooling rate of the steel sheet center portion is reduced to 0. controlled cooling at a cooling rate to be 5 to 20 ° C. / s,該制allowed to cool from your cooling stop temperature T4 (° C.) to room temperature, characterized in that the amount of hydrogen of the steel sheet below 0.65 ppm, resistance hydrogen Pulling with excellent induced cracking and ductile fracture characteristics Strength 760MPa grade above manufacturing method of the high strength steel sheet.
In place of the steel component,
C: 0.02-0.10%,
Si: 0.01 to 0.6%,
Mn: 1.5 to 2.5%
P: 0.015% or less,
S: 0.003% or less, Ni: 0.1-2.0%,
Mo: 0.15-0.60%,
Nb: 0.001 to 0.10%,
Ti: 0.005 to 0.030%,
Al: 0.06% or less,
N: 0.0001 to 0.006%,
The method for producing a high-strength steel sheet having a tensile strength of 760 MPa or more and excellent in hydrogen-induced cracking resistance and ductile fracture characteristics according to claim 1, comprising: Furthermore, in mass%,
B: 0.0001 to 0.005%,
V: 0.001 to 0.10%,
Cu: 0.01 to 1.0%,
Cr: 0.01 to 0.8%
Zr: 0.0001 to 0.005%,
Ta: 0.0001 to 0.005%,
Ca: 0.0001 to 0.01%,
REM: 0.0001 to 0.01%,
Mg: 0.0001 to 0.006%
High tensile strength of 760 MPa class or more excellent in hydrogen-induced crack resistance and ductile fracture characteristics according to any one of claims 1 to 3, characterized by containing at least one of A method of manufacturing a steel sheet.   The reheating temperature of the steel slab is 1000 to 1250 ° C, and the tensile strength of 760 MPa class excellent in hydrogen-induced crack resistance and ductile fracture characteristics according to any one of claims 2 to 4. The manufacturing method of the above high strength steel plate.   6. The tensile strength of 760 MPa excellent in hydrogen-induced crack resistance and ductile fracture characteristics according to claim 1, wherein after the controlled cooling is stopped, the steel plates are superposed and allowed to cool. A manufacturing method for high-strength steel sheets of grades or better.   The method for producing a high-strength steel sheet having a tensile strength of 760 MPa class or more excellent in hydrogen-induced crack resistance and ductile fracture characteristics according to claim 6, wherein the laminated steel sheet is covered with a heat insulating cover and allowed to cool. .   A tensile strength of 760 MPa class excellent in hydrogen-induced crack resistance and ductile fracture characteristics, characterized by being piped using the high-strength steel sheet produced by the method according to any one of claims 1 to 7. The manufacturing method of the above high strength welded steel pipe.   In the pipe forming process, the high-strength steel sheet is formed into a tubular shape in the UO process, the ends are submerged arc welded using a welding wire and a firing type flux or a melt type flux, and then the pipe is expanded. The method for producing a high-strength welded steel pipe having a tensile strength of 760 MPa or more and excellent in hydrogen-induced crack resistance and ductile fracture characteristics according to claim 8. % By mass
C: 0.01 to 0.12%,
Si: 0.3% or less,
Mn: 1.2-2.4%
Ni: 4.0 to 8.5%,
Cr + Mo + V: 3.0-5.0%,
Ti: 0.005 to 0.15%,
The hydrogen-induced crack resistance and ductile fracture according to claim 9, wherein submerged arc welding is performed using a welding wire containing Al: 0.02% or less, the balance being iron and inevitable impurities. A method for producing a high-strength welded steel pipe with excellent tensile strength of 760 MPa class or higher. 11. The hydrogen-induced crack resistance and ductile fracture according to claim 9, wherein the heat input per 1 mm thickness of the submerged arc welding is 0.13 to 0.25 kJ / mm ^2. A method for producing a high-strength welded steel pipe with excellent tensile strength of 760 MPa class or higher.

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