Title of invention: High-strength steel sheet
Technical field
[0001]
　The present invention is a high-strength steel plate, specifically, a high-strength steel plate having a tensile strength of 1200 MPa or more and excellent in seizure curability and weldability, which is suitable for structural members such as automobiles mainly used by being pressed. It is about.
　The present application claims priority based on Japanese Patent Application No. 2018-141226 filed in Japan on July 27, 2018, the contents of which are incorporated herein by reference.
Background technology
[0002]
　In recent years, in order to protect the global environment, it has been required to improve the fuel efficiency of automobiles, and in order to reduce the weight and ensure safety of automobile steel sheets, further increase in strength is required. When the strength of the steel sheet is increased, the ductility is generally lowered, which makes cold press forming difficult. Therefore, there is a demand for a material that is relatively soft during molding and is easy to mold, and has high strength after molding, that is, a material having excellent seizure curability.
[0003]
　The material having excellent baking curability as used herein is a material having a high amount of baking hardening and high strength after baking hardening.
[0004]
　The baking hardening is performed by diffusing penetrating elements (carbon and nitrogen) into dislocations formed by press molding (hereinafter, also referred to as “pre-strain”) during coating baking at 150 ° C. to 200 ° C. to fix the dislocations. It is a strain aging phenomenon that occurs.
[0005]
　As shown in Non-Patent Document 1, the amount of baking hardened depends on the amount of penetrating elements in solid solution, that is, the amount of solid solution carbon. Therefore, the amount of baking hardening is higher in martensite, which has a large amount of carbon that can be solid-solved than ferrite, which has a small amount of carbon that can be solid-solved. In connection with this, for example, Patent Document 1 discloses a high-strength steel plate mainly composed of bainite and martensite, and in the high-strength steel plate disclosed in Patent Document 1, a predetermined treatment is applied to the steel material. The baking hardness is improved by increasing the dislocation density. Considering these points, it is considered that even with the same martensite, the amount of baking hardening is increased by increasing the concentration of added carbon.
[0006]
　On the other hand, if too much carbon or alloying element is added, the weldability generally deteriorates. Carbon equivalent (Ceq) is one of the indicators of weldability. This is a method of estimating weldability from the component ratio contained in the steel sheet. For example, Ceq is defined by the following equation according to the JIS standard. Here, the content (mass%) of each element is substituted for each element symbol in the formula.
　Ceq = C + Si / 24 + Mn / 6 + Ni / 40 + Cr / 5 + Mo / 4 + V / 14
[0007]
　However, although the above formula is suitable for evaluation of high carbon thick steel sheets used for building materials, it is said that it is not suitable for automobile steel sheets. Therefore, as shown in Non-Patent Document 2, Ono has proposed Ceq shown in the following equation.
　Ceq = C + Si / 90 + (Mn + Cr) / 100 + 1.5P + 3S
[0008]
　Generally, the higher the Ceq, the more difficult it is to weld. Therefore, in order to improve weldability, it is important to reduce the elements contained in the above formula. In conventional high-strength steel sheets for automobiles, an upper limit is set for the C content, and weldability is ensured by supplementing the strength with other alloying elements. Such a technique is disclosed in, for example, Patent Document 2. That is, weldability is ensured by reducing the concentration of added carbon. It is also important to ensure the characteristics after welding. For example, in a structure containing island-shaped martensite (MA) in the matrix phase, since MA is a structure harder than the matrix phase, it acts as an embrittlement phase and deteriorates toughness after welding.
[0009]
　As described above, from the viewpoint of the alloy component, it is difficult to achieve both seizure curability and weldability at the same time.
[0010]
　Further, in the high-strength steel plate described in Patent Document 1, as described above, not only martensite and bainite are mainly used, but also the seizure curability is improved by increasing the dislocation density. However, in general, steel having a high dislocation density causes thermal strain embrittlement as shown in Non-Patent Document 3, and therefore has poor weldability.
[0011]
　On the other hand, in the invention described in Patent Document 3, weldability is ensured by precipitating metal carbides by tempering the matrix phase as martensite, bainite or the like. However, in the invention described in Patent Document 3, since there is a tempering step, there is a problem that the solid solution carbon is reduced and the seizure curability is deteriorated.
[0012]
　As described above, it is difficult to achieve both seizure curability and weldability not only from the viewpoint of alloy components but also from the viewpoint of dislocation density.
Prior art literature
Patent documents
[0013]
Patent Document 1: Japanese Patent Application Laid-Open No. 2008-144233
Patent Document 2: Japanese Patent Application Laid-Open No. 3-180445
Patent Document 3: Japanese Patent Application Laid-Open No. 2007-308743
Non-patent literature
[0014]
Non-Patent Document 1: K.K. Nakaoka, et al. , "Strength, Ductility and Aging Properties of Continuously-Annealed Dual-Phase High-Strength Sheet Steels", Formable HSLA and Steel- Soc. of AIME, (1977) 126-141
Non-Patent Document 2: Moriaki Ono, "Spot Weldability of High-Strength Thin Steel Sheets for Automobiles", Welding Technology, 51 (3) (2003) 77-82
Non-Patent Document 3: Kunihiko Sato Et al., Proceedings of the Japan Shipbuilding Society, 142 (1977) 173-181
Outline of the invention
Problems to be solved by the invention
[0015]
　In order to meet the demand for higher strength in the future, it is necessary to increase the carbon concentration in order to ensure excellent seizure curability. However, as a result, there is a problem that Ceq increases and weldability deteriorates. Further, from the viewpoint of dislocation density, it is difficult to achieve both seizure curability and weldability.
[0016]
　Therefore, an object of the present invention is to provide a high-strength steel plate having high seizure curability and excellent weldability.
Means to solve problems
[0017]
　The present inventors have attempted to ensure the above-mentioned seizure curability and weldability by the following two approaches.
(1) Ceq is suppressed and weldability is ensured by appropriately controlling the alloy components.
(2) In order to secure an appropriate amount of solid solution carbon, seizure curability should be obtained by using martensite as the parent phase while quenching.
[0018]
　However, this alone did not provide the desired tensile strength after baking and hardening. As a result of detailed investigation, the deformed structure after baking and curing was non-uniform. Therefore, the present inventors found that the prestrain was non-uniform due to the difference in hardness in martensite, and therefore all martensite. Could not be used for baking cure, and it was considered that the baking cure would deteriorate. Then, the present inventors have found that this non-uniform hardness difference is caused by the microsegregation of Mn. In general, microsegregation is a phenomenon in which the concentration of alloying elements generated from solidification is unevenly distributed, and planes perpendicular to the plate thickness direction are connected in layers.
[0019]
　Therefore, the present inventors have found that the heat spreading process is controlled to suppress the microsegregation of Mn by making it a uniform structure, and the prestrain is made uniform, so that the seizure curability is greatly improved. It was. In addition, the uniform structure makes it difficult for MA to occur and improves weldability.
[0020]
　The high-strength steel sheets having excellent seizure curability and weldability of the present invention that have achieved the above object in this way are as follows.
(1) In terms of mass%,
　C: 0.05 to 0.15%,
　Si: 1.5% or less,
　Mn: 2.00 to 5.00%,
　P: 0.100% or less,
　S: 0.010 %
　Or less, Al: 0.001 to 2.000%,
　N: 0.010% or less
, the balance
　consists of Fe and impurities , and Ceq defined by the following formula (1) is less than 0.21. ,
　Contains martensite of 98% or more in area ratio, has a residual structure of 2% or less in area ratio, and has
　a two-dimensional homogeneous dispersion ratio S defined by the formula (2) of 0.85 or more and 1.20 or less. A
　high-strength steel plate having a tensile strength of 1200 MPa or more.
　Ceq = C + Si / 90 + (Mn + Cr) / 100 + 1.5P + 3S Equation (1)
　S = Sy 2 / Sx 2　　　　　　　　　　　　　　　　　　　　　　　　　Equation (2)
　Here, each element symbol in the equation (1) includes the content (mass%) of each element. ) Is substituted, and 0 is substituted when the element is not included, and Sx in the equation (2) is substituted. 2 is the dispersion value of the Mn concentration profile data in the plate width direction, and Sy 2 is the dispersion value of the Mn concentration profile data in the plate thickness direction.
(2) The high-strength steel plate according to (1), wherein the residual structure is composed of retained austenite when the residual structure is present.
(3) Further, according to (1) or (2), a total of 0.100% or 　less of one or two types having
　Ti: 0.100% or less and
Nb: 0.100% or less in mass%
is contained. High-strength steel sheet.
(4) Further, 　any of (1) to (3), which contains 1 or 2 types of
　Cu: 1.000% or less and
Ni: 1.000% or less in total in an amount of 1.000% or less in mass%. The high-strength steel sheet described in item 1.
(5) Further, in mass%,
　W: 0.005% or less,
　Ca: 0.005% or less,
　Mg: 0.005% or less
　Rare earth metal (REM): 0.010% or less
One or more types The high-strength steel sheet according to any one of (1) to (4), which contains 0.010% or less in total.
(6) The high-strength steel sheet according to any one of (1) to (5), further containing B: 0.0030% or less in mass%.
(7) The high-strength steel sheet according to any one of (1) to (6), further containing Cr: 1.000% or less in mass%.
Effect of the invention
[0021]
　According to the present invention, a high-strength steel plate having excellent weldability and high seizure curability by making the microsegregation of Mn a uniform structure in as-quenched martensite in which the alloy component is controlled, specifically, It is possible to provide a high-strength steel plate having a tensile strength of 1350 MPa after quenching and hardening. After pressing, it is baked during painting to increase its strength, so it is suitable as a structural field in fields such as automobiles.
Mode for carrying out the invention
[0022]

　The high-strength steel plate according to the embodiment of the present invention has a mass% of
　C: 0.05 to 0.15%,
　Si: 1.5% or less,
　Mn: 2.00 to 5.00%. ,
　P: 0.100% or less,
　S: 0.010% or less,
　Al: 0.001 to 2.000%,
　N: 0.010% or less
, and the balance
　consists of Fe and impurities. The Ceq defined in 1) is less than 0.21
　, contains martensite of 98% or more in area ratio, and the residual structure is 2% or less in area ratio,
　and is two-dimensional defined by the formula (2). The homogeneous dispersion ratio S is 0.85 or more and 1.20 or less, and the
　tensile strength is 1200 MPa or more.
　Ceq = C + Si / 90 + (Mn + Cr) / 100 + 1.5P + 3S Equation (1)
　S = Sy 2 / Sx 2　　　　　　　　　　　　　　　　　　　　　　　　　Equation (2)
　Here, each element symbol in the equation (1) includes the content (mass%) of each element. ) Is substituted, and 0 is substituted when the element is not included, and Sx 2 in the equation (2) is substituted. Is the dispersion value of the Mn concentration profile data in the plate width direction, and Sy 2 is the dispersion value of the Mn concentration profile data in the plate thickness direction.
[0023]
　First, the chemical composition of the high-strength steel sheet according to the embodiment of the present invention and the slab used for manufacturing the same will be described. In the following description, "%", which is a unit of the content of each element contained in the high-strength steel sheet and the slab, means "mass%" unless otherwise specified.
[0024]
(C: 0.05% to 0.15%)
　C has an effect of increasing the amount of solid solution carbon and enhancing the baking curability. In addition, it has the effect of enhancing hardenability and increasing strength by containing it in the martensite structure. If the C content is less than 0.05%, a sufficient solid solution carbon content cannot be secured, and the baking hardening amount decreases. Therefore, the C content is set to 0.05% or more, preferably 0.08% or more. On the other hand, if the C content exceeds 0.15%, a silicate having a low melting point is produced during welding, which affects the quality of the weld seam. In addition, the strength is too high to guarantee moldability. Therefore, the C content is 0.15% or less, preferably less than 0.13%, 0.12% or less, 0.11% or less, or 0.10% or less.
[0025]
(Si: 1.5% or less)
　Si is a solid solution strengthening element and has a role of suppressing cementite precipitation, which is a factor for lowering the strength. Therefore, it may be included in the high-strength steel sheet of the present invention. On the other hand, if the Si content exceeds 1.5%, the surface texture may deteriorate. Therefore, the Si content is 1.5% or less, preferably 1.2% or less. The lower limit of the Si content is not particularly limited, but the content may be 0.01% or more because it functions as a deoxidizer for molten steel.
[0026]
(Mn: 2.00% to 5.00%)
　Mn is an element for improving hardenability, and is an element necessary for forming a martensite structure without limiting the cooling rate. In order to effectively exert this effect, the Mn content is 2.00% or more, preferably 2.50% or more. However, the content of excess Mn is 5.00% or less, preferably 4.50% or less because the low temperature toughness is lowered due to the precipitation of MnS.
[0027]
(P: 0.100% or less)
　P is not an essential element and is contained as an impurity in steel, for example. From the viewpoint of weldability, the lower the P content, the better. In particular, when the P content exceeds 0.100%, the weldability is significantly reduced. Therefore, the P content is 0.100% or less, preferably 0.030% or less. Reducing the P content is costly, and attempts to reduce it to less than 0.0001% significantly increase the cost. Therefore, the P content may be 0.0001% or more. Further, since P contributes to the improvement of strength, the P content may be 0.0001% or more from such a viewpoint.
[0028]
(S: 0.010% or less)
　S is not an essential element and is contained as an impurity in steel, for example. From the viewpoint of weldability, the lower the S content, the better. The higher the S content, the higher the amount of MnS deposited and the lower the low temperature toughness. In particular, when the S content exceeds 0.010%, the weldability and low temperature toughness are significantly reduced. Therefore, the S content is 0.010% or less, preferably 0.003% or less. Reducing the S content is costly, and attempts to reduce it to less than 0.0001% significantly increase the cost. Therefore, the S content may be 0.0001% or more.
[0029]
(Al: 0.001% to 2.000%)
　Al has an effect on deoxidation. In order to effectively exert the above actions, the Al content is 0.001% or more, preferably 0.010% or more. On the other hand, when the Al content exceeds 2.000%, the weldability is lowered, oxide-based inclusions are increased, and the surface texture is deteriorated. Therefore, the Al content is 2.000% or less, preferably 1.000% or less.
[0030]
(N: 0.010% or less)
　N is not an essential element and is contained as an impurity in steel, for example. From the viewpoint of weldability, the lower the N content, the better. In particular, when the N content exceeds 0.010%, the weldability is significantly reduced. Therefore, the N content is 0.010% or less, preferably 0.006% or less. Reducing the N content is costly, and attempts to reduce it to less than 0.0001% significantly increase the cost. Therefore, the N content may be 0.0001% or more.
[0031]
　The basic composition of the high-strength steel sheet of the present invention and the slab used for its production is as described above. Further, the high-strength steel sheet of the present invention and the slab used for producing the same may contain the following optional elements, if necessary.
[0032]
(Ti: 0.100% or less, Nb: 0.100% or less)
　Ti and Nb contribute to the improvement of strength. Therefore, Ti, Nb or any combination thereof may be contained. In order to sufficiently obtain this effect, the content of Ti or Nb, or the total content of the combination of these two types, is preferably 0.003% or more. On the other hand, if the content of Ti or Nb or the total content of the combination of these two types exceeds 0.100%, hot rolling and cold rolling become difficult. Therefore, the Ti content or Nb content, or the total content of the combination of these two types is set to 0.100% or less. That is, the limiting range in the case of each component alone is Ti: 0.003% to 0.100% and Nb: 0.003% to 0.100%, and the total content when these are combined is also , 0.003 to 0.100% is preferable.
[0033]
(Cu: 1.000% or less, Ni: 1.000% or less)
　Cu and Ni contribute to the improvement of strength. Therefore, Cu, Ni or a combination thereof may be contained. In order to sufficiently obtain this effect, the content of Cu and Ni is preferably in the range of 0.005 to 1.000% when each component is used alone, and the total content when these two types are combined , 0.005% or more and 1.000% or less is preferably satisfied. On the other hand, if the contents of Cu and Ni, or the total content when these two types are combined, exceeds 1.000%, the effect of the above action is saturated and the cost becomes unnecessarily high. Therefore, the upper limit of the content of Cu and Ni, or the total content when these two types are combined is 1.000%. That is, Cu: 0.005% to 1.000% and Ni: 0.005% to 1.000%, and the total content when these are combined is 0.005 to 1.000%. It is preferable to have.
[0034]
(W: 0.005% or less, Ca: 0.005% or less, Mg: 0.005% or less, REM: 0.010% or less)
　W, Ca, Mg and REM contribute to fine dispersion of inclusions. , Increase toughness. Therefore, W, Ca, Mg or REM or any combination thereof may be contained. In order to sufficiently obtain this effect, the total content of W, Ca, Mg and REM, or any combination of two or more thereof is preferably 0.0003% or more. On the other hand, when the total content of W, Ca, Mg and REM exceeds 0.010%, the surface texture deteriorates. Therefore, the total content of W, Ca, Mg and REM is 0.010% or less. That is, W: 0.005% or less, Ca: 0.005% or less, Mg: 0.005% or less, REM: 0.010% or less, and the total content of any two or more of these is 0. It is preferably .0003 to 0.010%.
[0035]
　REM (rare earth metal) refers to a total of 17 elements of Sc, Y and lanthanoids, and "REM content" means the total content of these 17 elements. Lanthanoids are industrially added, for example in the form of misch metal.
[0036]
(B: 0.0030% or less)
　B is an element for improving hardenability and is an element useful for forming a martensite structure. B is preferably contained in an amount of 0.0001% (1 ppm) or more. However, if B is contained in excess of 0.0030% (30 ppm), excessive boron causes high temperature brittleness and may affect welding performance. Therefore, the B content is set to 0.0030% or less. It is preferably 0.0025% or less.
[0037]
(Cr: 1.000% or less)
　Cr is an element for improving hardenability and is an element useful for forming a martensite structure. Cr is preferably contained in an amount of 0.005% or more. However, if Cr is contained in excess of 1.000%, the welding performance may be affected. Therefore, the Cr content is set to 1.000% or less. It is preferably 0.500%.
[0038]
　In the high-strength steel sheet according to the present embodiment, the balance other than the above components is composed of Fe and impurities. Here, the impurity is a component that is mixed due to various factors in the manufacturing process, including raw materials such as ore and scrap, when the high-strength steel sheet is industrially manufactured, and relates to the present embodiment. It means a component that is not intentionally added to a high-strength steel plate.
[0039]
(Ceq is less than 0.21) The
　present embodiment is characterized in that the Ceq represented by the following formula (1) is set to less than a predetermined value in order to improve weldability. Thereby, weldability can be ensured. In order to further enhance such an effect, it is necessary to secure Ceq at less than 0.21. It is preferably 0.18 or less.
　Ceq = C + Si / 90 + (Mn + Cr) /100+1.5P+3S Equation (1)
　Here, the content (mass%) of each element is substituted for each element symbol in the equation (1), and when the element is not included, 0 is substituted.
[0040]
　Next, the structure of the high-strength steel sheet according to the embodiment of the present invention will be described. The organizational requirements will be described below, but% related to the organizational fraction means "area ratio".
[0041]
(Martensite: 98% or more) The
　present embodiment is characterized in that martensite is secured at an area ratio of 98% or more. As a result, sufficient solid solution carbon can be secured, and as a result, seizure curability can be enhanced. In order to further enhance such an effect, it is necessary to secure 98% or more of martensite, for example, 100%.
[0042]
　In the present invention, the area ratio of martensite is determined as follows. First, a sample is taken with the thickness cross section perpendicular to the rolling direction of the steel plate as the observation surface, the observation surface is polished, and the structure at the position 1/4 of the thickness of the steel plate is SEM-EBSD (electron) at a magnification of 5000 times. Observe with a scanning electron microscope with an electron backscatter diffraction device), analyze the image with a field of 100 μm × 100 μm to measure the area ratio of martensite, and the average of these measured values ​​in any 5 or more fields is the book. It is determined as the area ratio of maltensite in the invention.
[0043]
(Remaining structure: 2% or less)
　According to the present invention, the remaining structure other than martensite has an area ratio of 2% or less. In order to further improve the seizure curability of the high-strength steel sheet, it is preferably 0%. When a residual tissue is present, the residual tissue can include any tissue and is not particularly limited, but for example, it preferably contains retained austenite or is composed of retained austenite. Trace amounts of retained austenite may be unavoidable depending on the composition of the steel and the manufacturing method. However, such a small amount of retained austenite not only does not adversely affect the seizure curability, but also contributes to the improvement of ductility by the TRIP (Transformation Induced Plasticity) effect when deformed. Can be done. Therefore, the residual structure may contain retained austenite in an area ratio of 2% or less. However, in order to further enhance the seizure curability, the residual structure does not contain residual austenite and is preferably 0%.
[0044]
　In the present invention, the area ratio of retained austenite is determined by X-ray diffraction measurement. Specifically, the portion from the surface of the steel sheet to the 1/4 position of the thickness of the steel sheet is removed by mechanical polishing and chemical polishing, and MoKα rays are used as characteristic X-rays to 1/4 of the depth from the surface of the steel sheet. The X-ray diffraction intensity at the position is measured. Then, from the integrated intensity ratios of the diffraction peaks of the body-centered cubic lattice (bcc) phases (200) and (211) and the face-centered cubic lattice (fcc) phases (200), (220) and (311), the following is obtained. The area ratio of retained austenite is calculated using the formula of.
　Sγ = (I 200f + I 220f + I 311f ) / (I 200b + I 211b ) × 100 In the
　above formula, Sγ is the area ratio of retained austenite, and I 200f , I 220f and I 311f are the fcc phases (200) and ( I 311f , respectively). 220) and (311) diffraction peak intensities, I 200b and I 211b indicate the intensity of the bcc phase (200) and (211) diffraction peaks, respectively.
[0045]
(Two-dimensional homogeneous dispersion ratio S is 0.85 or more and 1.20 or less) The
　two-dimensional homogeneous dispersion ratio is an index for evaluating the microsegregation of alloying elements. The two-dimensional homogeneous dispersion ratio represented by S is measured as follows. The plate width direction is the x direction, the plate thickness direction is the y direction, and the surface of the steel plate whose rolling direction is the normal direction (that is, the cross section in the thickness direction of the steel plate) is adjusted so that it can be observed, and then mirror-polished and EPMA. (Electronic probe microanalyzer) From one side to the other along the thickness direction (y direction) of the steel plate in the range of 100 μm × 100 μm at the center of the steel plate in the thickness direction cross section of the steel plate The Mn concentration at 200 points is measured at 0.5 μm intervals. Further, the Mn concentration at 200 points is similarly measured from one side to the other side at intervals of 0.5 μm along the direction (x direction) perpendicular to the thickness direction of the measured steel sheet. The dispersion values ​​Sx 2 and Sy 2 are obtained from the Mn concentration profiles in the x and y directions . Using these values, S is calculated by the following equation (2).
　S = Sy 2 / Sx 2　　　　　　　　　　　　　　　　　　　　　　　　　equation (2)
　Here, Sx 2 is the dispersion value of the Mn concentration profile data in the plate width direction, and Sx 2 = (1/200) × Σ (AA i ) 2In the formula, A is the average value of Mn concentrations at 200 points in the x direction, and A i represents the i-th Mn concentration in the x direction (i = 1 to 200). Similarly, Sy 2 is a dispersion value of Mn concentration profile data in the plate thickness direction, and is represented by Sy 2 = (1/200) × Σ (BB i ) 2. In the formula, B is 200 in the y direction. It is the average value of the Mn concentration of the points, and B i represents the i-th Mn concentration in the y direction (i = 1 to 200).
[0046]
　The present embodiment is characterized in that the Mn concentration distribution has a uniform structure (for example, a checkered pattern structure) due to relaxation of microsegregation. If this is less than 0.85, it cannot be said that the structure is sufficiently uniform, and the seizure curability is low. In addition, MA is generated and the weldability is not good. Therefore, S is required to be 0.85 or more. It is preferably 0.90 or more, more preferably 0.95 or more. On the other hand, as described above, when microsegregation is not controlled, a surface having a high concentration of Mn and a surface having a low concentration of Mn are connected in layers in the plate thickness direction, and this can be homogenized in the plate thickness direction and the plate width direction. is important. On the contrary, if the surface having a high concentration of Mn and the surface having a low concentration of Mn are connected in a layer in the plate thickness direction, the homogenization is not achieved. That is, the reciprocal of the lower limit value of S is the upper limit value. Therefore, S is set to 1.20 or less. It is preferably 1.15 or less, more preferably 1.10 or less.
[0047]
　Next, the mechanical properties of the present invention will be described.
[0048]
(Tensile strength: 1200 MPa or more)
　According to the high-strength steel plate of the present invention having the above composition and structure, high tensile strength, specifically 1200 MPa or more, can be achieved. Here, the reason why the tensile strength is set to 1200 MPa or more is to satisfy the demand for weight reduction of the automobile body. The tensile strength is preferably 1300 MPa or more, more preferably 1400 MPa or more.
[0049]
　According to the high-strength steel sheet of the present invention, it is possible to achieve excellent seizure curability. More specifically, according to the high-strength steel plate of the present invention, the stress when the test piece heat-treated at 170 ° C. for 20 minutes after applying the 2% prestrain is re-tensioned, and the stress when the 2% prestrain is applied. It is possible to achieve the baking hardening amount BH such that the value obtained by subtracting the above value is 130 MPa or more, preferably 150 MPa or more. If the BH value is less than 130 MPa, it is difficult to mold and the strength after baking and curing is low, so that it cannot be said that the baking curability is excellent. Further, according to the high-strength steel plate of the present invention, baking hardening is performed so that the stress when the test piece heat-treated at 170 ° C. for 20 minutes after applying 2% prestrain is re-tensioned is 1350 MPa or more, preferably 1400 MPa or more. Later tensile strength BHTS can be achieved. If the BHTS value is less than 1350 MPa, the strength after baking and curing is similarly low, so that it cannot be said that the baking curability is excellent.
[0050]

　Next, a preferable method for manufacturing a high-strength steel sheet according to the present embodiment will be described.
[0051]
　The following description is intended to illustrate a characteristic method for manufacturing the high-strength steel plate of the present invention, and the high-strength steel plate of the present invention is manufactured by a manufacturing method as described below. It is not intended to be limited to.
[0052]
　A preferred method for producing a high-strength steel plate of the present invention is a step of casting molten steel having the chemical composition described above to form a
　slab, and rough rolling in which the slab is roughly rolled in a temperature range of 1050 ° C. or higher and 1250 ° C. or lower. In the process, the rough rolling includes reverse rolling with a reduction rate of 30% or less per pass performed an even number of times in 2 passes or more and 16 passes or less, and the reduction rate difference between the two passes when making one round trip. Is 20% or less, the reduction rate of even times in one round trip is 5% or more higher than the reduction rate of odd times, and the rough rolling process is held for 5 seconds or more after the
　rough rolling. A finish rolling process in which finish rolling is performed in a temperature range of ° C. or higher and 1050 ° C. or lower, wherein the finish rolling is performed on four or more continuous rolling stands, the rolling reduction of the first stand is 15% or more, and finish rolling. A finish rolling process in
　which the rolled steel sheet is wound in a temperature range of 400 ° C. or less, a cold rolling process in which the
　obtained hot-rolled steel sheet is cold-rolled at a reduction rate of 15% or more and 45% or less, and a cold-rolled steel sheet obtained. Is heated at an average heating rate of 10 ° C./sec or more, held for 10 to 1000 seconds in a temperature range of Ac 3 or more and 1000 ° C. or less, and then cooled to 70 ° C. or less at an average cooling rate of 10 ° C./sec or more. step, and
　is characterized in that it comprises a skin-pass rolling process to skin pass rolling with the resulting steel sheet of 0.5% to 2.5% or less of reduction ratio. Hereinafter, each step will be described.
[0053]
(Slab Forming Step)
　First, a molten steel having the chemical composition of the high-strength steel sheet according to the present invention described above is cast to form a slab to be subjected to rough rolling. The casting method may be a normal casting method, and a continuous casting method, an ingot forming method, or the like can be adopted, but the continuous casting method is preferable from the viewpoint of productivity.
[0054]
(Rough Rolling Step) It
　is preferable to heat the slab to a solution temperature range of 1000 ° C. or higher and 1300 ° C. or lower before rough rolling. The heating holding time is not particularly specified, but it is preferable to hold the slab at the heating temperature for 30 minutes or more in order to bring the slab to a predetermined temperature. The heating holding time is preferably 10 hours or less, more preferably 5 hours or less, in order to suppress excessive scale loss. If the temperature of the slab after casting is 1050 ° C. or higher and 1250 ° C. or lower, the slab may be subjected to rough rolling as it is without being heated and held in the temperature range, and may be directly fed or rolled directly.
[0055]
　Next, by roughly rolling the slab by reverse rolling, the Mn segregated portion in the slab formed during solidification in the slab forming step is made into a uniform structure without forming a plate-shaped segregated portion extending in one direction. can do. The formation of the Mn concentration distribution having such a uniform structure will be described in more detail. First, on the surface of the slab cut perpendicular to the surface of the slab before the start of rough rolling, it can be observed that alloying elements such as Mn are concentrated in a comb-like form. Specifically, in the cut surface of the slab before rough rolling, the portion where the alloying element such as Mn is linearly concentrated is substantially perpendicular to the surface of the slab from both surfaces of the slab toward the inside. It is in a state where multiple lines are lined up.
[0056]
　On the other hand, in rough rolling, the surface of the slab is stretched in the rolling traveling direction for each rolling pass. The rolling traveling direction is the direction in which the slab advances with respect to the rolling roll. Then, as the surface of the slab is stretched in the traveling direction of rolling in this way, the Mn segregated portion growing inward from the surface of the slab is in a state of being inclined in the traveling direction of the slab for each rolling pass. Be made. In other words, rolling has a function of slightly tilting the Mn segregated portion extending in a comb shape toward the inside of the slab in the direction of rolling progress.
[0057]
　Here, in the case of so-called unidirectional rolling in which the traveling direction of the slab in each pass of rough rolling is always the same direction, the Mn segregated portion heads in the same direction for each pass while maintaining its own substantially straight state. The slope gradually increases. Then, at the end of the rough rolling, the Mn segregation portion is in a substantially parallel posture with respect to the surface of the slab while maintaining a substantially straight state, and a flat microsegregation is formed.
[0058]
　On the other hand, in the case of reverse rolling in which the traveling directions of the slabs in each rough rolling pass are alternately opposite, the Mn segregated portion inclined in the direction of the immediately preceding pass is in the opposite direction in the next pass. Receives the force to tilt. In this case, the Mn segregated portion has a bent shape. Therefore, in the reverse rolling, the Mn segregated portions are alternately bent in a zigzag shape by repeatedly performing the passes in the opposite directions alternately.
[0059]
　When a plurality of zigzag shapes that are alternately bent in this way are lined up, the plate-like microsegregation disappears and the Mn concentration distribution becomes uniformly intricate. When the Mn concentration distribution is ideally homogenized, the Mn concentration distribution appears in a substantially checkered pattern. The "checkerboard pattern" (Ichimatsu pattern) is a kind of lattice pattern, and is a pattern in which substantially squares (or approximately rectangles) having different colors are arranged alternately. In the present invention, a structure in which the Mn concentration distribution appears in a checkered pattern is referred to as a checkered structure. By adopting a uniform structure in which the two-dimensional homogeneous dispersion ratio S is 0.85 or more and 1.20 or less, Mn is more easily diffused by the heat treatment in the subsequent process, and a hot-rolled steel sheet having a more uniform Mn concentration can be obtained. Obtainable. Since the Mn concentration distribution is uniformly intricate throughout the steel sheet by the above reverse rolling, such a uniform structure is obtained not only in the sheet thickness cross section parallel to the rolling direction but also in the sheet whose rolling direction is normal. It is similarly formed in the thick cross section.
[0060]
　If the rough rolling temperature range is less than 1050 ° C., it becomes difficult to complete rolling at 850 ° C. or higher in the final pass of rough rolling, resulting in poor shape. Therefore, the rough rolling temperature range is preferably 1050 ° C. or higher. More preferably, it is 1100 ° C. or higher. If the rough rolling temperature range exceeds 1250 ° C., scale loss increases and there is a concern that slab cracking may occur. Therefore, the rough rolling temperature range is preferably 1250 ° C. or lower.
[0061]
　If the reduction rate per pass in rough rolling exceeds 30%, the shear stress during rolling becomes large and the Mn segregated portion becomes non-uniform. Therefore, the rolling reduction rate per pass in rough rolling is set to 30% or less. The smaller the rolling reduction, the smaller the shear strain during rolling, and a uniform structure can be obtained. Therefore, the lower limit of the rolling ratio is not particularly set, but 10% or more is preferable from the viewpoint of productivity.
[0062]
　In order to make the Mn concentration distribution a uniform structure, reverse rolling is preferably 2 passes or more, more preferably 4 passes or more. However, if it is applied in excess of 16 passes, it becomes difficult to secure a sufficient finish rolling temperature, so the number is 16 passes or less. Further, it is desirable that each pass in which the traveling directions are opposite to each other is performed the same number of times, that is, the total number of passes is an even number. However, in a general rough rolling line, the entry side and the exit side of rough rolling are located on opposite sides of the roll. For this reason, the number of passes (rolling) in the direction from the entry side to the exit side of rough rolling increases once. Then, in the final pass (rolling), the Mn segregated portion becomes a flat shape, and it becomes difficult to form a uniform structure. When rough rolling is performed on such a hot rolling line, it is preferable to leave a gap between rolls in the final pass and omit rolling.
[0063]
　In reverse rolling, if there is a difference in rolling reduction between two passes included in one reciprocating rolling, shape defects are likely to occur, and the Mn segregated portion becomes non-uniform, so that a uniform structure cannot be obtained. Therefore, during rough rolling, the rolling reduction difference between the two passes included in one round trip of reverse rolling is set to 20% or less. It is preferably 10% or less.
[0064]
　As will be described later, in order to miniaturize the recrystallized structure, tandem multi-stage rolling in finish rolling is effective, but tandem rolling tends to form flat microsegregation. In order to utilize tandem multi-step rolling, it is necessary to make the even-numbered rolling reduction in reverse rolling larger than the odd-numbered rolling, and control the microsegregation formed in the subsequent tandem rolling. The effect becomes remarkable when the reduction rate of even-numbered times (return route) is 5% or more higher than the reduction rate of odd-numbered times (outward route) in one round trip of reverse rolling. Therefore, it is preferable that the even-numbered rolling reduction rate is 5% or more higher than the odd-numbered rolling reduction rate in one round trip of reverse rolling.
[0065]
　In order to make the complex structure of Mn produced by reverse rolling in rough rolling uniform by austenite grain boundary movement, it is preferable to hold it for 5 seconds or more from rough rolling to finish rolling.
[0066]
(Finish rolling process)
　After reverse rolling in rough rolling, in order to narrow the interval of the Mn segregation zone caused by the dendrite secondary arm by increasing the rolling reduction in tandem rolling in finish rolling, finish rolling is performed in 4 steps. It is preferably carried out on one or more continuous rolling stands. If the finish rolling temperature is less than 850 ° C, recrystallization does not occur sufficiently and the structure is stretched in the rolling direction, and a plate-like structure due to the stretched structure is generated in the subsequent process. Therefore, the finish rolling temperature is 850 ° C. The above is preferable. More preferably, it is 900 ° C. or higher. On the other hand, when the finish rolling temperature exceeds 1050 ° C., it becomes difficult to generate fine recrystallized grains of austenite, it becomes difficult to segregate Mn at the grain boundaries, and the Mn segregation zone tends to become flat. Therefore, the finish rolling temperature is preferably 1050 ° C. or lower. If the temperature is appropriate, the rough-rolled steel sheet may be heated after the rough-rolling step and before the finish-rolling step, if necessary. Further, when the reduction ratio of the first stand for finish rolling is set to 15% or more, a large amount of recrystallized grains are generated, and Mn is easily dispersed uniformly by the subsequent grain boundary movement. As described above, by limiting not only the rough rolling process but also the finish rolling process, microsegregation of flat Mn can be suppressed. The "finish rolling temperature" means the surface temperature of the steel sheet from the start of finish rolling to the end of finish rolling. When finish rolling is performed so that the finish rolling temperature is within the above range, the so-called finish rolling start temperature (steel plate temperature in the first pass of finish rolling) and the finish rolling end temperature (in the last pass of finish rolling). The steel plate temperature) is also within the range of the finish rolling temperature described above.
[0067]
　If the take-up temperature exceeds 400 ° C., the surface texture deteriorates due to internal oxidation, so the take-up temperature is preferably 400 ° C. or lower. When the steel sheet structure has a homogeneous structure of martensite or bainite, it is easy to form a homogeneous structure by annealing, so that the winding temperature is more preferably 300 ° C. or lower.
[0068]
(Cold Rolling Step) The
　hot-rolled steel sheet obtained in the finish rolling step is pickled and then subjected to cold rolling to obtain a cold-rolled steel sheet. In order to maintain the martensite lath, the reduction rate is preferably 15% or more and 45% or less. If the rolling reduction in the cold rolling process exceeds 45%, the fine lath of martensite cannot be maintained and Mn is less likely to be segregated at the grain boundaries. The belt stretches. In such a flat layered Mn segregation zone, the dispersion of Mn is non-uniform, so that the two-dimensional homogeneous dispersion ratio of Mn is lower than the above-mentioned specified value. The pickling may be a normal pickling.
[0069]
(Annealing step)
　The steel sheet obtained through the above cold rolling step is annealed. For heating at the annealing temperature, the temperature is raised at an average heating rate of 10 ° C./sec or more, and heating is held for 10 to 1000 seconds in a temperature range of Ac 3 or more and 1000 ° C. or less. This temperature range and annealing time are for austenite transformation of the entire surface of the steel sheet. When the holding temperature exceeds 1000 ° C. or the annealing time exceeds 1000 seconds, the austenite particle size becomes coarse and martensite having a large lath width is formed, resulting in a decrease in toughness. Therefore, the annealing temperature is Ac 3 or more and 1000 ° C. or less, and the annealing time is 10 to 1000 seconds.
[0070]
　The Ac 3 points are calculated by the following formula. Substitute the mass% of the element for the element symbol in the following formula. Substitute 0% by mass for elements that do not contain it.
　Ac 3 = 881-335 x C + 22 x Si-24 x Mn-17 x Ni-1 x Cr-27 x Cu
[0071]
　After maintaining the annealing temperature, cooling is performed at an average cooling rate of 10 ° C./sec or higher. In order to freeze the tissue and efficiently induce martensitic transformation, the cooling rate should be high. However, if the temperature is lower than 10 ° C./sec, martensite is not sufficiently produced and the desired tissue cannot be controlled. Therefore, the temperature is set to 10 ° C./sec or higher.
[0072]
　The cooling stop temperature is 70 ° C. or lower. This is because martensite is generated as it is hardened on the entire surface by cooling. If cooling is stopped above 70 ° C, tissues other than martensite may appear. Further, even when martensite is generated, precipitates such as spheroidized iron carbide may be generated by self-tempering. In such a case, the solid solution carbon is reduced and the baking curability is lowered. Therefore, the cooling stop temperature is 70 ° C. or lower, preferably 60 ° C. or lower.
[0073]
(Skin pass rolling process) After the
　annealing process, skin pass rolling (tempering rolling) is performed. This is necessary in order to work-harden soft martensite and uniformly insert dislocations due to prestrain when there is still a difference in hardness within martensite even if the structure is uniform. In addition, when retained austenite remains, it has a role of increasing the martensite fraction by performing martensitic transformation by plastic working-induced transformation. This effect is not achieved by skin pass rolling at a rolling reduction of less than 0.5%. Therefore, the reduction rate is set to 0.5% or more. However, since it becomes difficult to control the plate thickness, it is preferable that the upper limit is 2.5%. More preferably, the reduction rate is 1.0% or less.
[0074]
　In this way, the high-strength steel sheet according to the embodiment of the present invention can be manufactured.
[0075]
　It should be noted that all of the above embodiments merely show examples of embodiment in carrying out the present invention, and the technical scope of the present invention should not be construed in a limited manner by these. That is, the present invention can be implemented in various forms without departing from the technical idea or its main features.
Example
[0076]
　Next, examples of the present invention will be described. The conditions in the examples are one condition example adopted for confirming the feasibility and effect of the present invention, and the present invention is not limited to this one condition example. The present invention can adopt various conditions as long as the gist of the present invention is not deviated and the object of the present invention is achieved.
[0077]
　A slab having the chemical composition shown in Table 1 was produced, the slab was heated to 1300 ° C. for 1 hour, and then rough-rolled and finish-rolled under the conditions shown in Table 2 to obtain a hot-rolled steel sheet. Then, the hot-rolled steel sheet was pickled and cold-rolled at the reduction ratio shown in Table 2 to obtain a cold-rolled steel sheet. Subsequently, annealing and skin pass rolling were performed under the conditions shown in Table 2. Each temperature shown in Table 2 is the surface temperature of the steel sheet. Further, in Table 2, "difference in rolling rate between one round trip inner pass (return path-outward path)" indicates the difference in rolling rate between two passes included in one round trip rolling in reverse rolling. In each example, reverse rolling including a plurality of reciprocating passes was performed, but the rolling reduction difference between the reciprocating passes was the same for all the reciprocating passes. For example, Example No. In 1, the "number of rough rolling passes" was 8, and the "difference in rolling rate between one round-trip inner pass (return route-outward route)" was 5%, as shown in the table. This is Example No. In No. 1, reverse rolling of 4 round trips was carried out, and it means that the return rolling reduction rate was 5% larger than the outward rolling reduction rate in all 4 round trips.
[0078]
　Ac 3 in Table 2 was calculated by the formula shown below. The mass% of the element was substituted for the element symbol in the following formula. For elements not contained, 0% by mass was substituted.
　Ac 3 = 881-335 x C + 22 x Si-24 x Mn-17 x Ni-1 x Cr-27 x Cu
[0079]
[table 1]

[0080]
[Table 2-1]

[0081]
[Table 2-2]

[0082]
　The area ratios of martensite and retained austenite were determined from the obtained cold-rolled steel sheet by SEM-EBSD and X-ray diffraction method.
[0083]
　In particular, the area ratio of martensite was determined as follows. First, a sample is taken with the thickness cross section perpendicular to the rolling direction of the steel sheet as the observation surface, the observation surface is polished, and the structure at 1/4 of the thickness of the steel sheet is observed by SEM-EBSD at a magnification of 5000 times. Then, the area ratio of martensite was measured by image analysis in a field of 100 μm × 100 μm, and the average of these measured values ​​in any of the five visual fields was determined as the area ratio of martensite. The area ratio of retained austenite was determined by X-ray diffraction measurement. Specifically, the portion from the surface of the steel sheet to the 1/4 position of the thickness of the steel sheet is removed by mechanical polishing and chemical polishing, and MoKα rays are used as characteristic X-rays to 1/4 of the depth from the surface of the steel sheet. The X-ray diffraction intensity at the position was measured. Then, from the integrated intensity ratios of the diffraction peaks of the body-centered cubic lattice (bcc) phases (200) and (211) and the face-centered cubic lattice (fcc) phases (200), (220) and (311), the following is obtained. The area ratio of retained austenite was calculated using the formula of.
　Sγ = (I 200f + I 220f + I 311f ) / (I 200b + I 211b ) × 100 In the
　above formula, Sγ is the area ratio of retained austenite, I 200f , I 220f and I 311f.Indicates the intensity of the diffraction peaks of the fcc phase (200), (220) and (311) , respectively, and I 200b and I 211b indicate the intensity of the diffraction peaks of the bcc phase (200) and (211), respectively.
[0084]
　Further, the two-dimensional homogeneous dispersion ratio represented by S was determined by an EMPA apparatus.
[0085]
　Further, the tensile strength TS, the elongation at break EL, the amount of baking hardening BH, and the tensile strength BHTS after baking hardening of the obtained cold-rolled steel sheet were measured. In the measurement of tensile strength TS, breaking elongation EL, seizure hardening amount BH, and tensile strength BHTS after seizure hardening, JIS No. 5 tensile test pieces with the direction perpendicular to the rolling direction as the longitudinal direction were collected and conformed to JIS Z 2241. Then, a tensile test was conducted. The seizure hardening amount BH is a value obtained by subtracting the stress at the time of applying the 2% prestrain from the stress when the test piece heat-treated at 170 ° C. for 20 minutes after applying the 2% prestrain is re-tensioned. The tensile strength BHTS after baking hardening is the stress when the test piece heat-treated at 170 ° C. for 20 minutes after applying 2% prestrain is re-tensioned. In order to satisfy the demand for weight reduction of the automobile body, the tensile strength is 1200 MPa or more, preferably 1300 MPa or more, and more preferably 1400 MPa or more. Further, since it is easy to mold, the elongation is preferably 5% or more. Further, for BH, if it is less than 130 MPa, it is difficult to mold and the strength after molding becomes low, so that 130 MPa or more is required to have excellent seizure curability. More preferably, it is 150 MPa or more. For BHTS, 1350 MPa or more is required to improve the collision performance by baking hardening. More preferably, it is 1400 MPa or more.
[0086]
　As an evaluation of weldability, test pieces were collected in accordance with JIS Z 3137, the same steel plates were spot welded to each other, and a cross tensile test was performed. Specifically, when the cross tensile test was performed on the welded material under the conditions that the electrode DR 6 mm-40R, the welding time was 15 cycles / 60 Hz, the pressing force was 400 kgf, and the nugget diameter was 6 mm by changing the current value, the base metal was used. The case of breakage was judged as pass (GOOD), and the case of breakage of the nugget was judged as fail (BAD).
[0087]
[Table 3]

[0088]
　[Evaluation Results] As
　shown in Table 3, in Examples 1, 3, 5, 6, 9, 13, 16, 20, 24, 27 and 28, excellent tensile strength, seizure curability and weldability are obtained. Was done. In each case, the tensile strength was 1200 MPa or more, the BH was 130 MPa or more, the BHTS was 1350 MPa or more, and the base metal was broken in the cross tensile test.
[0089]
　On the other hand, in Comparative Example 2, since there was no skin pass rolling, retained austenite remained and BH was low. In Comparative Example 4, since the S content was too large, Ceq was high and weldability was poor. In Comparative Example 7, since the annealing temperature was too low, a ferrite structure appeared and a sufficient martensite structure could not be obtained, and as a result, TS, BH and BHTS were low. In Comparative Example 8, since the annealing time was too short, the entire surface did not have a martensite structure, and TS, BH and BHTS were also low. In Comparative Example 10, since the average cooling rate in the annealing step was too slow, the entire surface did not have a martensite structure, and TS, BH and BHTS were low. In Comparative Example 11, since the C content was too low, the solid solution carbon content was reduced, and TS, BH and BHTS were low. In Comparative Example 12, the weldability was poor because the P content was too high. In Comparative Example 14, since the difference in rolling reduction between the two passes during one round trip in the rough rolling process was large, the Mn concentration distribution was not a uniform structure, the BH was low, and the weldability was poor. In Comparative Example 15, since the even-numbered rolling reduction rate in one round trip in the rough rolling process was smaller than the odd-numbered rolling reduction rate, the structure did not have a uniform Mn concentration distribution, the BH was low, and the weldability was poor. .. In Comparative Example 17, since the number of reverse rolling passes in the rough rolling step was an odd number, the structure did not have a uniform Mn concentration distribution, the BH was low, and the weldability was poor.
[0090]
　In Comparative Example 18, since the cooling stop temperature in the annealing step was high, structures other than martensite appeared, and iron carbides were precipitated to reduce the amount of solute carbon, so that the BH was low. In Comparative Example 19, TS, BH and BHTS were low because the Mn content was too low. In Comparative Example 21, since the reduction rate of reverse rolling in the rough rolling step was high, the structure did not have a uniform Mn concentration distribution, the BH was low, and the weldability was poor. In Comparative Example 22, the time from rough rolling to finish rolling was too short, the Mn concentration distribution became flat, the BH was low, and the weldability was poor. In Comparative Example 23, since the C content was too high, the area ratio of retained austenite (γ) was high, the BH was low, the Ceq was high, and the weldability was poor. In Comparative Example 25, since the number of rolling stands for finish rolling was small, the Mn concentration distribution became flat, the BH and BHTS were low, and the weldability was poor. In Comparative Example 26, the cold spreading ratio was high, the Mn concentration distribution was extended in the direction perpendicular to the plate thickness and became flat, the BH and BHTS were low, and the weldability was poor. In Comparative Example 29, the rolling reduction of the first stand for finish rolling was small, the Mn concentration distribution was flattened, the BH was low, and the weldability was poor. In Comparative Example 30, since the finish rolling temperature (finish rolling start temperature in Table 2) was too high, the Mn concentration distribution became flat, the BH was low, and the weldability was poor. In Comparative Example 31, the weldability was poor because the Al content was too high. In Comparative Example 32, the weldability was poor because the N content was too high. In Comparative Example 33, the weldability was poor because the Ceq was too high.
Industrial applicability
[0091]
　The high-strength steel plate having excellent seizure curability and weldability of the present invention can be used as a raw plate for structural materials of automobiles, particularly in the field of the automobile industry.
The scope of the claims
[Claim 1]
　By mass%,
　C: 0.05 to 0.15%,
　Si: 1.5% or less,
　Mn: 2.00 to 5.00%,
　P: 0.100% or less,
　S: 0.010% or less,
　Al: 0.001 to 2.000%,
　N: 0.010% or less
, the balance
　consists of Fe and impurities, the Ceq defined by the following formula (1) is less than 0.21, and the
　area ratio. Contains 98% or more of martensite, the residual structure is 2% or less in area ratio, and the
　two-dimensional homogeneous dispersion ratio S defined by the following formula (2) is 0.85 or more and 1.20 or less.
　A high-strength steel plate having a tensile strength of 1200 MPa or more.
　Ceq = C + Si / 90 + (Mn + Cr) / 100 + 1.5P + 3S Equation (1)
　S = Sy 2 / Sx 2　　Equation (2)
　Here, each element symbol in the equation (1) includes the content (mass%) of each element. ) Is substituted, 0 is substituted when the element is not included, Sx 2 in the equation (2) is the dispersion value of the Mn concentration profile data in the plate width direction, and Sy 2Is the dispersion value of the Mn concentration profile data in the plate thickness direction.
[Claim 2]
　The high-strength steel plate according to claim 1, wherein the residual structure is composed of retained austenite when the residual structure is present.
[Claim 3]
The high-strength steel sheet according to claim 1 or 2, further containing 0.100% or less in total of one or two types having
　Ti: 0.100% or less and
　Nb: 0.100% or less 　in mass%
.
[Claim 4]
　Further, according to 　any one of claims 1 to 3, a total of 1.000% or less of one or two types having
　Cu: 1.000% or less and
Ni: 1.000% or less in mass% is contained. High-strength steel plate.
[Claim 5]
　Furthermore, in mass%,
　W: 0.005% or less,
　Ca: 0.005% or less,
　Mg: 0.005% or less
　Rare earth metal (REM): 0.010% or less
One type or two or more types in total The high-strength steel sheet according to any one of claims 1 to 4, which contains 0.010% or less.
[Claim 6]
　The high-strength steel sheet according to any one of claims 1 to 5, further comprising B: 0.0030% or less in mass%.
[Claim 7]
　The high-strength steel sheet according to any one of claims 1 to 6, further comprising Cr: 1.000% or less in mass%.