JP2013108167A - Method of producing high strength steel plate of tensile strength of 950 mpa or greater, excellent in weldability and delayed fracture resistance - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a method of producing a high strength steel plate of a tensile strength of 950 MPa or greater, excellent in weldability and delayed fracture resistance relative to prior art steel.SOLUTION: The method of producing a high strength steel plate of a tensile strength of 950 MPa or greater, excellent in weldability and delayed fracture resistance, comprises: heating a steel of a composition comprising by mass, C: 0.03 to 0.25%, Si: 0.01 to 0.8%, Mn: 0.5 to 2%, P: 0.010% or less, S: 0.003% or less, Al: 0.005 to 0.1%, and N: 0.0005 to 0.008% with the balance being Fe and inevitable impurities, wherein a weld crack sensitivity index Pcm of 0.26% or lower, to an Actransformation point or higher; hot-rolling the steel at a cumulative rolling reduction of not higher than 80% in an un-crystallization temperature range; completing the hot rolling at an Artransformation point or higher; subsequently cooling the steel from the Artransformation point or higher down to a temperature of 250°C or lower at a cooling speed of at least 10°C/s; thereafter reheating the steel at a mean temperature increase rate of at least 1°C/s; and subjecting the steel to a tempering treatment with the highest achieved temperature kept in a range of 100 to 400°C.

Description

  The present invention relates to a method for producing a high-tensile steel plate excellent in weldability and delayed fracture resistance, and particularly relates to a suitable method for producing a high-tensile steel plate having a tensile strength of 950 MPa or more.

  In recent years, in the field of using steel materials such as construction industrial machines, tanks, penstocks, line pipes, etc., against the backdrop of increasing the size and weight of structures, the strength of steel materials to be used has been increasing and the amount of steel materials used has rapidly increased. It has increased.

  In response to such an increase in demand, Patent Documents 1, 2, and 3 propose 780 MPa class high-tensile thick steel plates that omit the tempering treatment.

  Furthermore, in order to meet the demand for higher strength, Patent Document 4 proposes a high-tensile steel plate having a yield strength of 885 MPa or more that omits the tempering treatment.

  Patent Documents 5 and 6 propose a method for producing a steel sheet having excellent hydrogen embrittlement resistance using various techniques such as fine dispersion of carbides.

JP 2007-277622 A JP 2007-277623 A JP 2009-263774 A JP 2011-012315 A JP 2006-206942 A JP 2007-009324 A

  The high-strength steel sheets that omit the tempering treatment described in Patent Documents 1, 2, and 3 have a tensile strength of about 780 MPa, but the cooling stop temperature is in the range of room temperature to 350 ° C., and non-uniform carbides are generated by self-annealing. It tends to occur and is insufficient from the viewpoint of delayed fracture resistance. In addition, it is necessary to reduce C in order to adjust the as-quenched strength, and it is necessary to add a large amount of alloy in order to ensure necessary characteristics, which increases the manufacturing cost from the viewpoint of alloy saving.

  A high strength steel sheet with a yield strength of 885 MPa or more that omits the tempering process described in Patent Document 4 is specified to have a cooling stop temperature of 200 to 400 ° C. and then cooled at 20 ° C./min or less. Similar to 1, 2, and 3, non-uniform carbides are likely to be generated by self-annealing, which is insufficient from the viewpoint of delayed fracture resistance. Moreover, since the strength is secured by auto-aging, it is difficult to say that the low temperature toughness is sufficient.

  Moreover, when trying to obtain a high-strength steel sheet having a tensile strength of 780 MPa or more by the methods described in Patent Documents 5 and 6, the tempering temperature is high, the productivity is inferior, and many alloys need to be added. There is a problem that weldability deteriorates.

  The present invention has been made in view of such circumstances, and an object of the present invention is to provide a method for producing a high-tensile steel sheet having a tensile strength of 950 MPa or more and superior in weldability and delayed fracture resistance to conventional steel materials.

  As a result of intensive research on a steel sheet having a strength of 950 MPa or more, weldability, and delayed fracture resistance, the present inventors have specified the rate of temperature rise and the tempering temperature at the center in the thickness direction of the steel during tempering. This effectively suppresses the formation of carbides, suppresses the accumulation of diffusible hydrogen in the carbides, and further reduces the amount of diffusible hydrogen in the steel by dehydrogenation during heating, even after tempering. It has been found that high strength steel sheets with better weldability and delayed fracture resistance than conventional steels can be obtained by making the alloy strength comparable to the strength during quenching and enabling alloy-saving design. .

  The present invention has been made on the basis of the above-described findings and further studies. The gist of the present invention is as follows.

[1] By mass%, C: 0.03 to 0.25%, Si: 0.01 to 0.8%, Mn: 0.5 to 2%, P: 0.010% or less, S: 0.00. 003% or less, Al: 0.005 to 0.1%, N: 0.0005 to 0.008%, the weld crack sensitivity index Pcm shown in the following formula (1) is 0.26% or less, A steel having a composition composed of Fe and unavoidable impurities in the balance is heated to the Ac ₃ transformation point or higher, and hot rolling is performed so that the cumulative reduction rate in the non-recrystallization temperature range is 80% or less, and the Ar ₃ transformation point. The hot rolling is completed as described above. Subsequently, the steel is cooled from the Ar ₃ transformation point or higher to a temperature of 250 ° C. or lower at a cooling rate of 10 ° C./s or higher and then reheated at an average heating rate of 1 ° C./s or higher. Weldability and delayed fracture characterized by performing a tempering treatment in which the ultimate temperature is in the range of 100 to 400 ° C. Excellent tensile strength 950MPa or more of the method for manufacturing the high-tensile steel sheet properties.

  [2] In addition to the steel, by mass%, Mo: 0.01-1%, Nb: 0.001-0.1%, V: 0.001-0.5%, Ti: 0.03% The weldability according to [1] above, which contains at least one selected from Cu: 2% or less, Ni: 4% or less, Cr: 2% or less, and W: 2% or less A method for producing a high-tensile steel sheet having a tensile strength of 950 MPa or more and excellent in delayed fracture resistance.

  [3] One or more kinds selected from B: 0.0003% to 0.003%, Ca: 0.01% or less, and REM: 0.02% or less are further added to the steel by mass%. The method for producing a high-tensile steel sheet having a tensile strength of 950 MPa or more and having excellent weldability and delayed fracture resistance as described in [1] or [2] above.

  [4] The above [1] to [1], wherein the metal structure is a structure mainly composed of a martensite phase, the equivalent circle diameter of the carbide at the lath interface of the martensite phase is 100 nm or less, and the carbide coverage is 20% or less. [3] A method for producing a high-tensile steel sheet having a tensile strength of 950 MPa or more and excellent in weldability and delayed fracture resistance.

  By using the present invention, a high-strength steel having a high tensile strength of 950 MPa or more and excellent in weldability and delayed fracture resistance can be stably produced at low cost.

  The reasons for limiting the respective constituent requirements of the present invention will be described below.

1. About component composition First, the reason which prescribed | regulated the component composition of the steel of this invention is demonstrated. In addition, all component% means the mass%.

C: 0.03-0.25%
C is an indispensable element for obtaining the strength required for structural steel, but if it is added less than 0.03%, sufficient strength cannot be obtained, and a large amount of alloying element is required, so that weldability is required. Decreases. If the addition exceeds 0.25%, the amount of martensite generated in the weld heat-affected zone increases and the toughness decreases, so the C content is in the range of 0.03 to 0.25%. Preferably it is 0.12 to 0.23% of range. More preferably, it is 0.12 to 0.22% of range.

Si: 0.01 to 0.8%
Si is added for deoxidation, but if less than 0.01% is added, the deoxidation effect is not sufficient, and if added over 0.8%, the toughness of the base metal and the weld heat-affected zone significantly decreases. At the same time, the weldability is remarkably lowered, so the Si content is in the range of 0.01% to 0.8%. Preferably it is 0.05 to 0.6% of range. More preferably, it is 0.2 to 0.5% of range.

Mn: 0.5-2%
Mn is added from the viewpoint of securing the strength of the base metal. However, if it is added in an amount of less than 0.5%, the effect is not sufficient. Since the toughness is remarkably lowered, the amount of Mn is set in the range of 0.5 to 2%. Preferably it is 0.6 to 1.6% of range. More preferably, it is 0.6 to 1.5% of range.

P: 0.010% or less When P exceeds 0.010%, the toughness of the base metal and the weld heat-affected zone is remarkably reduced, so the P content is 0.010% or less.

S: 0.003% or less If S is contained in excess of 0.003%, the toughness of the base metal and the weld heat-affected zone is remarkably reduced. Therefore, the amount of S is made 0.003% or less. Preferably it is 0.0015% or less.

Al: 0.005 to 0.1%
Al is added in order to sufficiently deoxidize molten steel, but if it is added less than 0.005%, the deoxidation effect is not sufficient, and if added over 0.1%, the amount of Al that dissolves in the base metal And the base metal toughness is reduced, so the Al content is in the range of 0.005 to 0.1%. Preferably it is 0.01 to 0.06% of range. More preferably, it is 0.01 to 0.04% of range.

N: 0.0005 to 0.008%
N is added in order to refine the structure by forming a nitride with Ti or the like and to improve the toughness of the base material and the weld heat affected zone. However, if the addition is less than 0.0005%, the effect of refining the structure is not sufficient. On the other hand, if the addition exceeds 0.008%, the amount of N dissolved in the base material increases and the base material toughness is remarkably increased. Further, a coarse carbonitride is formed also in the weld heat affected zone and the toughness is lowered. Therefore, the N content is set in the range of 0.0005% to 0.008%. Preferably it is 0.0010 to 0.006% of range. More preferably, it is 0.0010 to 0.004% of range.

Pcm: 0.26% or less The weld cracking susceptibility index Pcm value is set to 0.26% or less in order to reduce the preheating temperature during welding and to make the steel material easy to construct. If the Pcm is too low, the required strength cannot be ensured, so the range is preferably 0.15 to 0.26%, more preferably 0.20 to 0.26%.
Here, Pcm is obtained by the following formula (1).

  The above is the basic chemical component of the present invention, and the balance consists of Fe and unavoidable impurities, but one kind selected from Mo, Nb, V, Ti, Cu, Ni, Cr, and W for the purpose of further increasing the strength. The above may be added as selective elements.

Mo: 0.01 to 1%
Mo is an element effective for increasing the strength of the base material, but if it is less than 0.01%, the effect is not sufficient, and if added over 1%, it causes an increase in hardness due to precipitation of alloy carbides, resulting in toughness. In order to reduce, when adding Mo, it is preferable to make Mo amount into the range of 0.01 to 1%. More preferably, it is 0.1 to 0.8% of range.

Nb: 0.001 to 0.1%
Nb is an element effective for strengthening steel, but if it is less than 0.001%, the effect is not sufficient, and if it exceeds 0.1%, the toughness of the base metal is significantly reduced. The Nb content is preferably in the range of 0.001 to 0.1%. More preferably, it is 0.005 to 0.05% of range.

V: 0.001 to 0.5%
V is effective in improving the strength and toughness of the base metal, and is effective for lowering the solid solution N by being precipitated as VN. However, if it is less than 0.001%, the effect is not sufficient. If added in excess of%, the toughness decreases due to precipitation of hard VC, so when adding V, the V content is preferably in the range of 0.001 to 0.5%. More preferably, it is 0.01 to 0.1% of range.

Ti: 0.03% or less Ti produces TiN at the time of rolling heating or welding, and effectively suppresses the austenite coarsening and improves the toughness of the base metal and the weld heat affected zone. If added in excess, Ti nitride becomes coarse and lowers the toughness of the base metal and the weld heat affected zone. Therefore, when adding Ti, the Ti content is preferably 0.03% or less. More preferably, it is 0.02% or less.

Cu: 2% or less Cu can improve the strength of the steel without impairing the low-temperature toughness, but if added over 2%, the steel plate surface is cracked during hot rolling. The amount is preferably 2% or less. More preferably, it is 1% or less.

Ni: 4% or less Ni is a beneficial element that improves the strength of the steel and the toughness of the heat affected zone of the steel, but if added over 4%, the effect is saturated and the economy is inferior, so Ni is added. In this case, the amount of Ni is preferably 4% or less. More preferably, it is 2% or less.

Cr: 2% or less Cr is an element effective for improving the strength and toughness, but if added over 2%, the weldability deteriorates. Therefore, when Cr is added, the Cr amount is 2% or less. It is preferable to do. More preferably, it is 0.1 to 1.2% of range.

W: 2% or less W is an element that has the effect of improving the strength, but if added over 2%, the weldability deteriorates, so when adding W, the amount of W is 2% or less. It is preferable that More preferably, it is 0.05 to 2% of range.

  In addition to the above composition, the high-tensile steel of the present invention may be added with one or more selected from B, Ca, and REM as selective elements for the purpose of further improving the material.

B: 0.0003 to 0.003%
B has the effect of suppressing the ferrite transformation from the grain boundary by segregating at the austenite grain boundary and improving the hardenability, but 0.0003% or more may be added in order to fully exhibit this effect. Preferably, if added over 0.003%, it precipitates as carbonitride, lowering the hardenability and lowering the toughness. Therefore, when adding B, the amount of B is 0.0003 to 0.003%. It is preferable to be in the range. More preferably, it is 0.0005 to 0.002% of range.

Ca: 0.01% or less Ca is an element useful for controlling the form of sulfide inclusions. However, if added over 0.01%, the cleanliness is lowered, so when adding Ca, the Ca content is preferably 0.01% or less. More preferably, it is 0.0005 to 0.0025% of range.

REM: 0.02% or less REM also has the effect of improving the quality of the material by forming oxides and sulfides in the steel, similar to Ca, but the effect is saturated even if added over 0.02%. Therefore, when REM is added, the amount of REM is preferably 0.02% or less. More preferably, it is 0.0005 to 0.0025% of range.

2. Regarding the metal structure, the reasons for limiting the metal structure in the present invention will be described.

  In order to increase the tensile strength of 950 MPa or more, the metal structure is a structure mainly composed of martensite. In the present invention, the structure mainly composed of martensite has a martensite + lower bainite structure fraction of 95% or more in order to achieve both strength and toughness. Less than 5% ferrite, upper bainite, residual γ, etc. are allowed.

  The microstructure fraction is quantified by etching with nital, taking 10 or more views of the structure photograph with a scanning electron microscope, and measuring the area ratio of martensite and lower bainite by image analysis.

The martensitic structure is tempered from the viewpoint of improving the weldability and delayed fracture resistance by reducing the amount of diffusible hydrogen in the steel as much as possible.
In addition, fine dispersion of carbide is extremely important for improving delayed fracture resistance. In particular, by reducing the equivalent circle diameter of the carbide generated at the lath interface of the martensite phase to 100 nm or less, the amount of carbide in the lath interface (hereinafter referred to as carbide coverage) is set to 20% or less. In addition, accumulation of diffusible hydrogen in the carbide can be suppressed, and the delayed fracture resistance can be improved. On the other hand, if the equivalent circle diameter of the carbide at the lath interface exceeds 100 nm or the carbide coverage at the lath interface exceeds 20%, the delayed fracture resistance deteriorates due to accumulation of diffusible hydrogen in the carbide.
In addition, about the measuring method of the circle equivalent diameter and carbide | carbonized_material coverage of the carbide | carbonized_material produced | generated in the lath interface of the martensite, after etching with nital, a structure | tissue photograph was image | photographed 10 or more views with the scanning electron microscope, and the photograph was used. Ask. For example, the equivalent circle diameter of carbide is converted into an equivalent circle diameter by analyzing the size of carbide generated at the interface of 50 or more laths by image analysis. In addition, the carbide coverage is measured by image analysis of the length (L _carbide ) of the carbide formed on 50 or more lath interfaces and the interface length (L _lath ) of the _lath by image analysis. The sum of the lengths along the interface is divided by the sum of the lath interface lengths and multiplied by 100.

3. About manufacturing conditions The manufacturing method of this invention is demonstrated below.

  In the present invention, the steel having the above-described composition is melted by a melting means such as a converter or an electric furnace, and is made into a steel material such as a slab by a conventional method such as a continuous casting method or an ingot-bundling method. However, it does not specify a method for melting or casting steel.

The steel slab having the composition described above for rolling conditions, heated to Ac ₃ transformation point or higher in a heating furnace. As a method of charging a steel slab into the heating furnace, a hot piece charging method in which the slab is charged into the heating furnace without being cooled below the Ar ₃ transformation point, or a slab once cooled is charged into the heating furnace. However, although there is a cold piece charging method in which reheating is performed to the Ac ₃ transformation point or higher, any method can be used in the present invention.

The reason why the steel is heated to the Ac ₃ transformation point or higher in the heating furnace is to make the steel uniform in one phase of the austenite structure, and the heating temperature is preferably 1100 ° C. or higher and 1250 ° C. or lower. In particular, when emphasizing toughness, the temperature is more preferably 1100 ° C. or more and 1200 ° C. or less.

Hot rolling, the cumulative reduction rate in the pre-recrystallization temperature region is 80% or less, and finished hot rolled at Ar ₃ transformation point or higher, the accelerated cooling start temperature is made to be than the Ar ₃ transformation point followed by .

The cumulative rolling reduction is 80% or less. However, when the cumulative rolling reduction in the non-recrystallization temperature region exceeds 80%, strain introduction by controlled rolling becomes significant, and the driving force for transformation increases, after rolling, or It causes hardenability degradation such as the formation of some ferrite during cooling, and it becomes difficult to achieve both high strength and high toughness of the base material of 950 MPa or more, so the cumulative rolling reduction is 80% or less. To do. In addition, since the austenite grain refinement effect cannot be obtained if the cumulative rolling reduction is less than 10%,
A preferable range of the cumulative rolling reduction is in the range of 10 to 80%.

  The non-recrystallization temperature range is a temperature range where recrystallization does not occur during rolling, and is 930 ° C. or less in the steel of the present invention. The cumulative rolling reduction is expressed by (original thickness−finished thickness) / original thickness × 100%.

Incidentally, Ar ₃ transformation point, a value that is calculated by the following equation (2).

As the Ac ₃ transformation point, a value calculated by the following formula (3) is used.

Cooling conditions after rolling After hot rolling is completed, forced cooling is performed from a temperature above the Ar ₃ transformation point to a temperature below 250 ° C. at a cooling rate of 10 ° C./s or more in order to ensure the base metal strength and toughness.

When cooling is started from the Ar ₃ transformation point or lower, a part of ferrite is generated, and it is difficult to achieve a base material strength of 950 MPa or higher. Therefore, it is necessary to start cooling from a temperature above the Ar ₃ transformation point.

  The reason for cooling until the cooling stop temperature is 250 ° C. or lower is to complete the transformation from austenite to martensite and strengthen the base material. In the present invention, in order to achieve both strength and toughness, the martensite + lower bainite structure fraction is set to 95% or more.

  The cooling rate during forced cooling is 10 ° C./s or more. This is because if it is less than 10 ° C./s, ferrite and pearlite are likely to be partially formed during cooling, and the desired strength and toughness cannot be secured stably.

  As a cooling method, techniques such as direct quenching and accelerated cooling are used, but the cooling method is not specified if a cooling rate of 10 ° C./s or more and a cooling stop temperature of 250 ° C. or less are obtained.

  In addition, a cooling rate is prescribed | regulated by the average cooling rate in 700-500 degreeC. This is because this temperature range is a temperature range in which a soft phase such as ferrite or pearlite is likely to appear, and in order to obtain a high-strength martensite structure, it is necessary to cool this temperature range quickly.

Tempering conditions Tempering is preferably performed using a heating device installed directly on the same production line as the rolling mill and the direct cooling or accelerated cooling device. This is because the time required from the rolling / cooling process to the tempering process can be shortened by the direct connection, thereby improving the productivity and reducing the heat energy.

  As the heating method for tempering, any method such as induction heating, current heating, infrared radiation heating, atmosphere heating, or the like may be used as long as the average temperature rising rate can be achieved and the upper and lower limits of the heating temperature can be managed.

  The temperature conditions for tempering were an average rate of temperature increase of 1 ° C./s or higher and an upper limit of the heating temperature of 400 ° C. By increasing the heating rate during heating to 1 ° C./s or higher, it is possible to suppress the coarsening of carbides at the lath interface that occurs at a slow heating rate such as in normal furnace heating. Moreover, by restrict | limiting the upper limit of heating temperature to 400 degreeC, the spreading | diffusion of C can be suppressed and the production | generation and coarsening of a carbide | carbonized_material can be suppressed effectively. On the other hand, when the heating temperature is less than 100 ° C, securing of stable strength and dehydrogenation are insufficient, and the lower limit of the heating temperature is set to 100 ° C or more.

  The above-described temperatures of the present invention are all temperatures at the center of the plate thickness unless otherwise specified, and are managed by calculation from the surface measured temperature. In addition, an average temperature increase rate is the value which divided the temperature increase amount required for reheating to reheating temperature (100-400 degreeC) after cooling by the time required for reheating.

  Moreover, the temperature rising process at the time of tempering is not limited as long as a predetermined average temperature rising rate can be obtained, and a linear temperature history may be taken or a temperature history that stays at an intermediate temperature may be taken.

  The holding time at the tempering temperature is preferably 60 s or less in order to prevent deterioration of toughness due to productivity / manufacturing cost and coarsening of precipitates.

  With respect to the cooling rate after tempering, the average cooling rate at the central portion of the plate thickness up to 100 ° C. or lower is 0.05 ° C./s or higher, 20 to prevent deterioration of toughness due to coarsening of precipitates during cooling. It is desirable that the temperature is not more than ° C / s.

  Twenty kinds of steel types A to T having chemical components shown in Table 1 were melted to cast a slab, heated in a heating furnace, and then rolled to obtain a steel plate. After rolling, it was subsequently quenched directly, and then tempered using a solenoid induction heating device.

  The average heating rate at the center of the plate thickness was controlled by the plate passing rate of the steel plate. In addition, when hold | maintaining at tempering temperature, it hold | maintained in the range of +/- 5 degreeC by reciprocating and heating a steel plate.

  The cooling after heating was air cooling. The temperature at the center of the plate thickness, such as the tempering temperature and the quenching temperature, was obtained by heat transfer calculation from the results of temperature measurement at the surface in succession by a radiation thermometer.

  Table 2 shows steel sheet production conditions, and yield strength, tensile strength, fracture surface transition temperature (vTrs), preheating temperature, resistance to delayed fracture safety, and martensite + lower bainite structure fraction in metal structure (% ), Equivalent circle diameter of carbide, and carbide coverage.

  Tensile tests are performed in accordance with JIS Z 2241. Yield strength and tensile strength are measured with a JIS No. 5 test piece when the plate thickness is 20 mm or less, and with a JIS No. 4 test piece taken from 1/4 part of the plate thickness when the plate thickness exceeds 20 mm. did.

  The toughness was evaluated by a fracture surface transition temperature (vTrs) obtained by a Charpy impact test using a test piece taken from 1/4 part of the plate thickness after collecting an impact test piece specified in JIS Z 2242.

  Weldability was evaluated by a preheating temperature necessary for preventing root cracking by performing coated arc welding with a heat input of 17 kJ / cm by a method prescribed in JIS Z 3158.

The delayed fracture resistance index is obtained by charging hydrogen so that the amount of diffusible hydrogen in the test piece is about 0.5 massppm by the cathodic hydrogen charging method using a smooth round bar test piece, and then galvanizing the surface of the test piece. Hydrogen is sealed by applying, and then a tensile test is performed at a strain rate of 1 × 10 ⁻⁶ / s to obtain a squeeze of the fractured test piece, while no hydrogen charge is performed at the same strain rate. A tensile test of the piece was also performed and evaluated according to the following formula.
Delayed fracture resistance index (%) = 100 x (X1 / X0)
Here, X0: Drawing of test specimen substantially not containing diffusible hydrogen (%)
X1: Test specimen containing diffusible hydrogen (%)
The target values for each characteristic were a yield stress of 885 MPa or more, a tensile strength of 950 MPa or more, vTrs ≦ −40 ° C., and a necessary preheating temperature of 75 ° C. or less. The target for the delayed fracture resistance index was 80% or more.

  In addition, the structure of the steel sheet was quantified by revealing the structure with a nital etchant in the vicinity of ¼ part of the plate thickness, performing 10 field observations with a scanning electron microscope, and evaluating the average value by image analysis.

  As shown in Table 2, Invention Example No. manufactured by the method of the present invention was used. 1-No. No. 12 has a component composition and production method within the scope of the present invention, and a steel sheet having good strength, weldability and delayed fracture resistance can be obtained.

  On the other hand, comparative example No. in which a component composition remove | deviates from this invention. 13-No. In No. 20, at least one of the properties of strength, toughness, preheating temperature, delayed fracture safety index, carbide equivalent circle diameter, and carbide coverage does not reach the target value. Further, the composition of the components is within the scope of the present invention, but the production conditions deviate from the present invention. 21-No. In No. 28, any one or more of the properties of strength, toughness, preheating temperature, delayed fracture safety index, carbide equivalent circle diameter, and carbide coverage does not reach the target value.

Claims (4)

In mass%, C: 0.03-0.25%, Si: 0.01-0.8%, Mn: 0.5-2%, P: 0.010% or less, S: 0.003% or less , Al: 0.005 to 0.1%, N: 0.0005 to 0.008%, the weld cracking sensitivity index Pcm shown in the following formula (1) is 0.26% or less, and the balance is Fe In addition, the steel having a component composition composed of inevitable impurities is heated to the Ac ₃ transformation point or higher, hot rolling is performed so that the cumulative reduction in the non-recrystallization temperature range is 80% or less, and the steel is heated above the Ar ₃ transformation point. After the hot rolling was completed, the steel sheet was cooled to a temperature of 250 ° C. or lower at a cooling rate of 10 ° C./s or higher from the Ar ₃ transformation point or higher, and then reheated at an average temperature rising rate of 1 ° C./s or higher. Weldability and delayed fracture resistance characterized by performing a tempering treatment in the range of 100 to 400 ° C. Excellent tensile strength of 950MPa or more of the method of manufacturing the high-tensile steel plate.

  In addition to the steel, in mass%, Mo: 0.01-1%, Nb: 0.001-0.1%, V: 0.001-0.5%, Ti: 0.03% or less, Cu 2 or less, Ni: 4% or less, Cr: 2% or less, and W: 2% or less. The weldability and delayed fracture resistance according to claim 1 A method for producing a high-tensile steel sheet having an excellent tensile strength of 950 MPa or more.   One or more selected from B: 0.0003% to 0.003%, Ca: 0.01% or less, and REM: 0.02% or less are further added to the steel in mass%. The method for producing a high-tensile steel sheet having a tensile strength of 950 MPa or more and excellent in weldability and delayed fracture resistance according to claim 1.   The metal structure is a structure mainly composed of a martensite phase, the equivalent circle diameter of carbide at the lath interface of the martensite phase is 100 nm or less, and the carbide coverage is 20% or less. A method for producing a high-tensile steel plate having a tensile strength of 950 MPa or more and excellent in weldability and delayed fracture resistance.