steel plate

Abstract

This steel plate has a prescribed chemical composition and has a microstructure containing, in terms of area ratio, 5.0% or less of ferrite and over 90.0% in total of martensite and tempered martensite, the balance being one or more of bainite, pearlite, and retained austenite. In a cross-section taken along the plate thickness direction, regions in which the Mn content is 1.1×[Mnave] or greater are present such that the number density of said regins is 5.0×10−4 / μm2 or greater and such that the spacing in relation to the nearest other region in which the Mn content is 1.1×[Mnave] or greater is, on average, 10.0 μm or less, where [Mnave] is the average of the Mn content in the entirety of the plate thickness direction. The steel plate has a tensile strength of 1,500 MPa or greater.

Description

[Technical Field]

[0001] The present invention relates to a steel sheet. This application claims priority based on Japanese Patent Application No. 2021-172425, filed on October 21, 2021, the contents of which are incorporated herein by reference. [Background technology]

[0002] With the highly specialized nature of industrial technology today, materials used in each field are required to have specialized and advanced performance. In particular, the demand for high-strength steel sheets for automobiles has increased significantly due to environmental concerns and the need to reduce vehicle weight and improve fuel efficiency. However, as strength increases, many metallic materials experience deterioration in their properties, particularly increased susceptibility to hydrogen embrittlement. It is known that steel components are particularly susceptible to hydrogen embrittlement when their tensile strength exceeds 1200 MPa. There have been cases of hydrogen embrittlement cracking in bolt steel, which has been used for high-strength applications prior to the automotive industry. Therefore, a fundamental solution to hydrogen embrittlement is strongly desired for high-strength steel sheets with tensile strengths of 1500 MPa or higher.

[0003] High-strength steel sheets with a tensile strength of 1500 MPa or more often have a microstructure that is primarily martensite or tempered martensite, but in such high-strength steel sheets, hydrogen that penetrates the steel segregates at the grain boundaries of martensite, embrittling the grain boundaries (reducing grain boundary strength) and causing cracks (hydrogen embrittlement). Because hydrogen penetration occurs even at room temperature, there is no method to completely prevent it, and a fundamental solution requires modifying the internal structure of the steel. Many proposals have been made so far regarding techniques for improving the hydrogen embrittlement resistance (sometimes referred to as hydrogen embrittlement resistance characteristics) of high-strength steel sheets (see, for example, Patent Documents 1 to 6).

[0004] Patent Document 1 describes an ultra-high strength steel sheet excellent in hydrogen embrittlement resistance and workability, which satisfies, in mass%, C: over 0.25 to 0.60%, Si: 1.0 to 3.0%, Mn: 1.0 to 3.5%, P: 0.15% or less, S: 0.02% or less, Al: 1.5% or less (excluding 0%), Mo: 1.0% or less (excluding 0%), Nb: 0.1% or less (excluding 0%), with the balance consisting of iron and unavoidable impurities, and which is subjected to a tensile work of 3%. The present invention discloses an ultra-high strength thin steel sheet with excellent hydrogen embrittlement resistance, characterized in that the metal structure after processing satisfies, in area ratios relative to the total structure, retained austenite structure: 1% or more, bainitic ferrite and martensite: 80% or more in total, and ferrite and pearlite: 9% or less in total (including 0%), while also satisfying an average axial ratio (major axis / minor axis) of the retained austenite grains: 5 or more, and having a tensile strength of 1180 MPa or more.

[0005] Patent Document 2 discloses a high-strength steel sheet having a tensile strength of 1500 MPa or more, which contains 1.0% or more of Si+Mn in its steel components, has a main phase structure in which ferrite and carbides form layers, and further, the lamellar structure in which the carbides have an aspect ratio of 10 or more and the spacing between the layers is 50 nm or less accounts for 65% or more in volume fraction relative to the entire structure, and further, the area fraction of carbides that form layers with ferrite and have an aspect ratio of 10 or more and that form angles of 25° or less with respect to the rolling direction is 75% or more, thereby providing a high-strength steel sheet with excellent bendability in the rolling direction and resistance to delayed fracture.

[0006] Patent Document 3 discloses a thin, ultra-high strength cold-rolled steel sheet with excellent bendability and delayed fracture resistance, which contains, by mass%, C: 0.15 to 0.30%, Si: 0.01 to 1.8%, Mn: 1.5 to 3.0%, P: 0.05% or less, S: 0.005% or less, Al: 0.005 to 0.05%, N: 0.005% or less, with the balance being Fe and unavoidable impurities, has a steel sheet surface soft portion that satisfies the relationship "hardness of the steel sheet surface soft portion / hardness of the steel sheet center portion ≦ 0.8", the proportion of the steel sheet surface soft portion to the sheet thickness is 0.10 to 0.30, the volume fraction of the steel sheet surface soft portion is tempered martensite 90% or more, the structure of the steel sheet center portion is tempered martensite, and the tensile strength is 1270 MPa or more.

[0007] Patent Document 4 discloses a cold-rolled steel sheet having a tensile strength of 1470 MPa or more and excellent bending workability and delayed fracture resistance, the cold-rolled steel sheet containing, by mass%, C: 0.15 to 0.20%, Si: 1.0 to 2.0%, Mn: 1.5 to 2.5%, P: 0.020% or less, S: 0.005% or less, Al: 0.01 to 0.05%, N: 0.005% or less, Ti: 0.1% or less, Nb: 0.1% or less, B: 5 to 30 ppm, with the balance being Fe and unavoidable impurities, and having a metallographic structure in which the volume fraction of the tempered martensite phase is 97% or more and the volume fraction of the retained austenite phase is less than 3%.

[0008] Patent Document 5 discloses an ultrahigh-strength steel sheet having a tensile strength of 1470 MPa or more that can exhibit excellent delayed fracture resistance even at cut edges, the ultrahigh-strength steel sheet having a composition containing, by mass%, 0.15 to 0.4% C, 0.5 to 3.0% Mn, and 0.001 to 0.10% Al, with the balance being iron and unavoidable impurities, with P, S, and N of the unavoidable impurities being limited to P: 0.1% or less, S: 0.01% or less, and N: 0.01% or less, respectively, and having a structure consisting of 90% or more martensite and 0.5% or more retained austenite, in terms of area fraction relative to the entire structure, in which 2% or more regions exist in which the local Mn concentration is 1.1 times or more the Mn content of the entire steel sheet, in terms of area fraction. The ultrahigh-strength steel sheet has a tensile strength of 1470 MPa or more.

[0009] Patent Document 6 discloses an ultra-high strength cold rolled steel sheet having excellent hydrogen embrittlement resistance and a tensile strength of 1300 MPa or more, which contains C: 0.150 to 0.300%, Si: 0.001 to 2.0%, Mn: 2.10 to 4.0%, P: 0.05% or less, S: 0.01% or less, N: 0.01% or less, Al: 0.001% to 1.0%, Ti: 0.001% to 0.10%, and B: 0.0001% to 0.010%, in which the values of the amount of solute B solB [mass%] and the prior austenite grain size Dγ [μm] satisfy the relationship solB·Dγ≧0.0010, the content of polygonal ferrite is 10% or less, the content of bainite is 30% or less, the content of retained austenite is 6% or less, and the content of tempered martensite is 60% or more, and the number density of Fe carbides in the tempered martensite is 1×10 6 / mm 2 The average dislocation density of the entire steel is 1.0 × 10 or more. 15 ~2.0×10 16 / m 2 The present invention discloses an ultra-high strength cold-rolled steel sheet having a steel structure with a grain size of 7.0 μm or less. It has been done. [Prior art documents] [Patent documents]

[0010] [Patent Document 1] Japanese Patent Application Publication No. 2006-207019 [Patent Document 2] Japanese Patent Application Publication No. 2010-138489 [Patent Document 3] Japanese Patent Publication No. 2011-179030 [Patent Document 4] Japanese Patent Publication No. 2010-215958 [Patent Document 5] Japanese Patent Application Publication No. 2016-153524 [Patent Document 6] Japanese Patent Application Publication No. 2016-050343 Summary of the Invention [Problem to be solved by the invention]

[0011] As mentioned above, several techniques have been proposed to improve the hydrogen embrittlement resistance of high-strength steel sheets. However, Patent Document 1 only discloses the hydrogen embrittlement resistance when a stress of 1000 MPa is applied, and does not provide any technical solutions for improving the hydrogen embrittlement resistance when a higher stress is applied. As mentioned above, hydrogen embrittlement occurs when hydrogen accumulates at grain boundaries, reducing the bond strength of the grain boundaries. Therefore, in order to prevent hydrogen embrittlement, it is effective to uniformly and finely disperse regions in the steel that have a stronger attractive interaction with H (hydrogen) than the prior austenite grain boundaries, thereby preventing H from accumulating on the prior γ grain boundaries. However, Patent Documents 1 to 6 do not consider methods for improving hydrogen embrittlement resistance from this perspective. In recent years, requirements for hydrogen embrittlement resistance have become even stricter, and Patent Documents 1 to 6 may not be able to meet such strict requirements. That is, conventionally, there has been room for improvement in hydrogen embrittlement resistance in high strength steel sheets having a microstructure mainly composed of martensite and tempered martensite. Therefore, an object of the present invention is to provide a steel sheet having high strength and excellent resistance to hydrogen embrittlement. [Means for solving the problem]

[0012] As mentioned above, hydrogen embrittlement is thought to be a crack that occurs at the grain boundary due to a decrease in the bonding strength of the grain boundary caused by the segregation of hydrogen in the steel at the grain boundary (mainly the prior austenite grain boundary in the case of a microstructure mainly consisting of martensite and tempered martensite). Therefore, the inventors investigated ways to prevent H from accumulating on the prior γ grain boundaries by uniformly and finely dispersing regions in the steel that have a stronger attractive interaction with H (hydrogen) than the prior austenite grain boundaries, and focused on a method that utilizes the attractive interaction of Mn with H. As a result, the following findings were obtained: a) By dispersing regions with a higher Mn content than the average Mn content at predetermined intervals, it is possible to suppress the accumulation of H at the prior austenite grain boundaries. b) In high-strength steel sheets having a microstructure mainly composed of martensite and tempered martensite, in order to disperse regions with a higher-than-average Mn content, it is important to control the dispersion state of cementite by optimizing the rolling conditions and coiling conditions in the hot rolling process, and then to control the annealing conditions and the cooling conditions after annealing.

[0013] The present invention has been made in light of the above findings. [1] A steel sheet according to one embodiment of the present invention contains, by mass%, C: 0.150 to 0.400%, Si: 0.01 to 2.00%, Mn: 0.8 to 2.0%, P: 0.0001 to 0.0200%, S: 0.0001 to 0.0200%, Al: 0.001 to 1.000%, N: 0.0001 to 0.0200%, O: 0.0001 to 0.0200%, Co: 0 to 0.500%, Ni: 0 to 1.000%, Mo: 0 to 1.000%, Cr: 0 to 2.000%, Ti: 0 to 0.500%, B: 0 to 0.0100%, Nb: 0 to 0.500%, V: 0 to 0.500%, and Cu: 0 to 0.500%. a chemical composition consisting of W: 0 to 0.100%, Ta: 0 to 0.100%, Mg: 0 to 0.050%, Ca: 0 to 0.050%, Y: 0 to 0.050%, Zr: 0 to 0.050%, La: 0 to 0.050%, Ce: 0 to 0.050%, Sn: 0 to 0.050%, Sb: 0 to 0.050%, As: 0 to 0.050%, and the balance: Fe and impurities; and a microstructure consisting of, in area ratio, ferrite: 5.0% or less, martensite and tempered martensite: more than 90.0% in total, and the balance: one or more of bainite, pearlite, and retained austenite; Cross section parallel to the plate thickness direction and rolling direction In this case, when the average Mn content of the entire plate thickness direction is [Mnave], the region where the Mn content is 1.1 × [Mnave] or more has a number density of 5.0 × 10 -4 pieces / μm 2 or more, and the distance to the nearest region where the Mn content is 1.1 × [Mnave] or more is 10.0 μm or less on average, and the tensile strength is 1500 MPa or more. [2] The steel sheet according to [1], wherein the chemical composition is Co: 0.01 to 0.500%, Ni: 0.01 to 1.000%, Mo: 0.01 to 1.000%, Cr: 0.001 to 2.000%, Ti: 0.001 to 0.500%, B: 0.0001 to 0.0100%, Nb: 0.001 to 0.500%, V: 0.001 to 0.500%, Cu: 0.001 to 0.500%, W: 0.001 to 0.100%, Ta: 0.001 to 0.100%, Mg: 0.001and As: 0.001 to 0.050%. [3] The steel sheet according to [1] or [2] may have a coating layer containing zinc, aluminum, magnesium or an alloy thereof on the surface. [Effects of the Invention]

[0014] According to the above-described aspects of the present invention, a steel sheet having high strength and excellent hydrogen embrittlement resistance can be provided. [Brief explanation of the drawings]

[0015] [Figure 1] 1 is a graph showing the influence of the average interval of the regions satisfying 1.1×[Mnave] and the number density of the regions satisfying 1.1×[Mnave] on hydrogen embrittlement resistance. DETAILED DESCRIPTION OF THE INVENTION

[0016] Hereinafter, a steel sheet according to one embodiment of the present invention (a steel sheet according to this embodiment) will be described. The steel sheet according to this embodiment has a predetermined chemical composition, and its microstructure is composed of, in area ratio, ferrite: 5.0% or less, martensite and tempered martensite: more than 90.0% in total, and the remainder: one or more of bainite, pearlite, and retained austenite, and in a cross section in the sheet thickness direction, when the average Mn content of the entire sheet thickness direction is [Mnave], the number density of a region where the Mn content is 1.1 × [Mnave] or more is 5.0 × 10 -4 pieces / μm 2 or more, and the distance to the nearest region where the Mn content is 1.1 × [Mnave] or more is 10.0 μm or less on average, and the tensile strength is 1500 MPa or more.

[0017] <Chemical composition> First, the range of the content of each element constituting the chemical composition of the steel sheet according to this embodiment will be described. Hereinafter, "%" relating to the content of an element means "% by mass." Furthermore, ranges indicated with "to" include both end values as the lower and upper limits.

[0018] C: 0.150 to 0.400% C is an element that is effective for increasing tensile strength inexpensively. If the C content is less than 0.150%, the target tensile strength cannot be obtained, and the fatigue properties of the welded portion deteriorate. For this reason, the C content is set to 0.150% or more. The C content may also be 0.160% or more, 0.180% or more, or 0.200% or more. On the other hand, if the C content exceeds 0.400%, hydrogen embrittlement resistance and weldability deteriorate. Therefore, the C content is set to 0.400% or less. The C content may be 0.350% or less, 0.300% or less, or 0.250% or less.

[0019] Si: 0.01 to 2.00% Silicon acts as a deoxidizer and is an element that affects the morphology of carbides and retained austenite after heat treatment. If the Si content is less than 0.01%, it becomes difficult to suppress the formation of coarse oxides. These coarse oxides become the starting point for cracks, which propagate within the steel material, deteriorating hydrogen embrittlement resistance. For this reason, the Si content is set to 0.01% or more. The Si content may be 0.05% or more, 0.10% or more, or 0.30% or more. On the other hand, if the Si content exceeds 2.00%, local ductility may decrease and hydrogen embrittlement resistance may deteriorate. Therefore, the Si content is set to 2.00% or less. The Si content may be 1.80% or less, 1.60% or less, or 1.40% or less.

[0020] Mn: 0.8 to 2.0% Mn is an element that is effective in increasing the strength of steel sheets. If the Mn content is less than 0.8%, the effect is not sufficiently obtained. Therefore, the Mn content is set to 0.8% or more. The Mn content may be 1.0% or more, or 1.2% or more. On the other hand, if the Mn content exceeds 2.0%, Mn not only promotes co-segregation with P and S, but may also deteriorate corrosion resistance and hydrogen embrittlement resistance. Therefore, the Mn content is set to 2.0% or less. The Mn content may also be 1.9% or less, or 1.8% or less.

[0021] P: 0.0001 to 0.0200% P is an element that strongly segregates at ferrite grain boundaries and promotes grain boundary embrittlement. If the P content exceeds 0.0200%, hydrogen embrittlement resistance is significantly reduced due to grain boundary embrittlement. Therefore, the P content is set to 0.0200% or less. The P content may also be 0.0180% or less, 0.0150% or less, or 0.0120% or less. The lower the P content, the better. However, if the P content is less than 0.0001%, the time required for refining increases, resulting in a significant increase in costs. For this reason, the P content is set to 0.0001% or more. The P content may be 0.0005% or more, 0.0010% or more, or 0.0020% or more.

[0022] S: 0.0001 to 0.0200% S is an element that forms non-metallic inclusions such as MnS in steel. If the S content exceeds 0.0200%, the formation of non-metallic inclusions that serve as crack initiation sites during cold working becomes significant. In this case, cracks are initiated from the non-metallic inclusions, and these cracks propagate through the steel, deteriorating its hydrogen embrittlement resistance. For this reason, the S content is set to 0.0200% or less. The S content may also be 0.0180% or less, 0.0150% or less, or 0.0120% or less. The lower the S content, the better. However, if the S content is less than 0.0001%, the time required for refining increases, resulting in a significant increase in costs. For this reason, the S content is set to 0.0001% or more. The S content may be 0.0005% or more, 0.0010% or more, or 0.0020% or more.

[0023] Al: 0.001 to 1.000% Al acts as a deoxidizer for steel and stabilizes ferrite. If the Al content is less than 0.001%, the effect is not sufficiently obtained. Therefore, the Al content is set to 0.001% or more. The Al content may be 0.005% or more, 0.010% or more, or 0.020% or more. On the other hand, if the Al content exceeds 1.000%, coarse Al oxides are formed. These coarse oxides become the starting points for cracks. Therefore, when coarse Al oxides are formed, cracks are generated in the coarse oxides, and these cracks propagate within the steel material, deteriorating the hydrogen embrittlement resistance. For this reason, the Al content is set to 1.000% or less. The Al content may also be 0.950% or less, 0.900% or less, or 0.800% or less.

[0024] N: 0.0001 to 0.0200% N is an element that forms coarse nitrides in steel sheets, reducing the hydrogen embrittlement resistance of the steel sheets, and is also an element that causes blowholes during welding. If the N content exceeds 0.0200%, hydrogen embrittlement resistance deteriorates and blowholes occur significantly. Therefore, the N content is set to 0.0200% or less. The N content may be 0.0180% or less, 0.0160% or less, or 0.0120% or less. On the other hand, if the N content is less than 0.0001%, the manufacturing cost increases significantly. Therefore, the N content is set to 0.0001% or more. The N content may be 0.0005% or more, 0.0010% or more, or 0.0020% or more.

[0025] O: 0.0001 to 0.0200% O is an element that forms oxides and deteriorates hydrogen embrittlement resistance. In particular, oxides are often present as inclusions, and when present on punched or cut edges, they form notch-like scratches or large dimples on the edge surface, which lead to stress concentration during heavy working and become the starting point for crack formation, resulting in a significant deterioration of workability. If the O content exceeds 0.0200%, the tendency for the above-mentioned deterioration of workability becomes significant. Therefore, the O content is set to 0.0200% or less. The O content may be 0.0180% or less, 0.0150% or less, or 0.0100% or less. A small O content is preferable. However, an O content of less than 0.0001% leads to excessive cost increases and is not economically preferable. For this reason, the O content is set to 0.0001% or more. The O content may be 0.0005% or more, 0.0010% or more, or 0.0015% or more.

[0026] The basic components of the chemical composition of the steel sheet according to the embodiment of the present invention are as described above. That is, the chemical composition of the steel sheet according to the present embodiment may include the above, with the balance being Fe and impurities. On the other hand, the chemical composition of the steel sheet according to the present embodiment may contain one or more of Co, Ni, Mo, Cr, Ti, B, Nb, V, Cu, W, Ta, Mg, Ca, Y, Zr, La, Ce, Sn, Sb, and As as optional components in place of a portion of the balance Fe, in order to improve various properties. These elements do not necessarily have to be contained, so the lower limit is 0%. Furthermore, even if the following elements are contained as impurities, the effects of the steel sheet according to this embodiment are not impaired.

[0027] Co: 0 to 0.500% Co is an element effective in controlling the morphology of carbides and increasing the strength of the steel sheet. Therefore, Co may be contained. To obtain a sufficient effect, the Co content is preferably 0.010% or more. The Co content may be 0.020% or more, 0.050% or more, or 0.100% or more. On the other hand, if the Co content exceeds 0.500%, coarse Co carbides precipitate. In this case, hydrogen embrittlement resistance may deteriorate. Therefore, the Co content is set to 0.500% or less. The Co content may be 0.450% or less, 0.400% or less, or 0.300% or less.

[0028] Ni: 0 to 1.000% Ni is an element effective in increasing the strength of steel sheets. Ni is also an element effective in improving wettability and accelerating alloying reactions. Therefore, Ni may be contained. To obtain the above effects, the Ni content is preferably 0.010% or more. The Ni content may be 0.020% or more, 0.050% or more, or 0.100% or more. On the other hand, if the Ni content exceeds 1.000%, hydrogen embrittlement resistance may decrease. Therefore, the Ni content is set to 1.000% or less. The Ni content may be 0.900% or less, 0.800% or less, or 0.600% or less.

[0029] Mo: 0 to 1.000% Mo is an element effective in increasing the strength of steel sheets. Mo also has the effect of suppressing ferrite transformation that occurs during heat treatment in continuous annealing equipment or continuous hot-dip galvanizing equipment. Therefore, Mo may be contained. To obtain the above effect, the Mo content is preferably 0.010% or more. The Mo content may be 0.020% or more, 0.050% or more, or 0.080% or more. On the other hand, if the Mo content exceeds 1.000%, the effect of suppressing ferrite transformation saturates. Therefore, the Mo content is set to 1.000% or less. The Mo content may be 0.900% or less, 0.800% or less, or 0.600% or less.

[0030] Cr: 0 to 2.000% Like Mn, Cr is an element that suppresses pearlite transformation and is effective in increasing the strength of steel. Therefore, Cr may be contained. To obtain the above effect, the Cr content is preferably 0.001% or more. The Cr content may be 0.005% or more, 0.010% or more, or 0.050% or more. On the other hand, if the Cr content exceeds 2.000%, coarse Cr carbides may form in the central segregation region, which may reduce hydrogen embrittlement resistance. Therefore, the Cr content is set to 2.000% or less. The Cr content may also be 1.800% or less, 1.500% or less, or 1.000% or less.

[0031] Ti: 0 to 0.500% Ti is an element that contributes to increasing the strength of steel sheets by strengthening precipitates, strengthening fine grains by inhibiting ferrite grain growth, and strengthening dislocations by inhibiting recrystallization. Therefore, Ti may be added. To obtain the above effects, the Ti content is preferably 0.001% or more. The Ti content may be 0.003% or more, 0.010% or more, or 0.050% or more. On the other hand, if the Ti content exceeds 0.500%, the precipitation of carbonitrides increases and hydrogen embrittlement resistance may deteriorate. Therefore, the Ti content is set to 0.500% or less. The Ti content may be 0.450% or less, 0.400% or less, or 0.300% or less.

[0032] B: 0 to 0.0100% B is an element that suppresses the formation of ferrite and pearlite during the cooling process from the austenite temperature range and promotes the formation of low-temperature transformation structures such as bainite or martensite. B is also an element that is beneficial for increasing the strength of steel. Therefore, B may be contained. To obtain the above effects, the B content is preferably 0.0001% or more. The B content may be 0.0003% or more, 0.0005% or more, or 0.0010% or more. On the other hand, if the B content exceeds 0.0100%, coarse B oxides are formed in the steel. These oxides become the starting points for voids during cold working, and the formation of coarse B oxides may deteriorate hydrogen embrittlement resistance. For this reason, the B content is set to 0.0100% or less. The B content may also be 0.0080% or less, 0.0060% or less, or 0.0050% or less.

[0033] Nb: 0 to 0.500% Nb, like Ti, is an element effective in controlling the morphology of carbides and is also effective in improving toughness by refining the structure. Therefore, Nb may be contained. To obtain the above effects, the Nb content is preferably 0.001% or more. The Nb content may be 0.002% or more, 0.010% or more, or 0.020% or more. On the other hand, if the Nb content exceeds 0.500%, the formation of coarse Nb carbides becomes significant. These coarse Nb carbides are prone to cracking, and the formation of coarse Nb carbides may deteriorate hydrogen embrittlement resistance. For this reason, the Nb content is set to 0.500% or less. The Nb content may also be 0.450% or less, 0.400% or less, or 0.300% or less.

[0034] V: 0 to 0.500% V is an element that contributes to increasing the strength of steel sheets by strengthening precipitates, strengthening fine grains by inhibiting ferrite grain growth, and strengthening dislocations by inhibiting recrystallization. Therefore, V may be added. To obtain the above effects, the V content is preferably 0.001% or more. The V content may be 0.002% or more, 0.010% or more, or 0.020% or more. On the other hand, if the V content exceeds 0.500%, the precipitation of carbonitrides increases and hydrogen embrittlement resistance may deteriorate. Therefore, the V content is set to 0.500% or less. The V content may be 0.450% or less, 0.400% or less, or 0.300% or less.

[0035] Cu: 0 to 0.500% Cu is an element effective in improving the strength of steel sheet. If the Cu content is less than 0.001%, a sufficient effect cannot be obtained. Therefore, to obtain the above effect, the Cu content is preferably 0.001% or more. The Cu content may be 0.002% or more, 0.010% or more, or 0.030% or more. On the other hand, if the Cu content exceeds 0.500%, hydrogen embrittlement resistance may deteriorate. Furthermore, if the Cu content is too high, the steel may become embrittled during hot rolling, making hot rolling impossible. For this reason, the Cu content is set to 0.500% or less. The Cu content may also be 0.450% or less, 0.400% or less, or 0.300% or less.

[0036] W: 0 to 0.100% W is an element effective in increasing the strength of steel sheet. W is also an element that forms precipitates and crystallized products. Precipitates and crystallized products containing W act as hydrogen trapping sites, so W is an element effective in improving hydrogen embrittlement resistance. Therefore, W may be contained. To obtain the above effects, the W content is preferably 0.001% or more. The W content may be 0.002% or more, 0.005% or more, or 0.010% or more. On the other hand, if the W content exceeds 0.100%, the formation of coarse W precipitates or crystallized particles becomes significant. These coarse W precipitates or crystallized particles are prone to cracking, and these cracks propagate within the steel material under low applied stress. Therefore, the formation of coarse W precipitates or crystallized particles may deteriorate hydrogen embrittlement resistance. For this reason, the W content is set to 0.100% or less. The W content may also be 0.080% or less, 0.060% or less, or 0.050% or less.

[0037] Ta: 0 to 0.100% Ta, like Nb, V, and W, is an element that is effective in controlling the morphology of carbides and increasing the strength of the steel sheet. Therefore, Ta may be contained. To obtain the above effects, the Ta content is preferably 0.001% or more. The Ta content may be 0.002% or more, 0.005% or more, or 0.010% or more. On the other hand, if the Ta content exceeds 0.100%, a large number of fine Ta carbides precipitate, and as the strength of the steel sheet increases, the ductility and bending resistance and hydrogen embrittlement resistance may decrease. Therefore, the Ta content is set to 0.100% or less. The Ta content may also be 0.080% or less, 0.060% or less, or 0.050% or less.

[0038] Mg: 0 to 0.050% Mg is an element that can control the morphology of sulfides even with a small amount of content. Therefore, Mg may be contained. To obtain the above effects, the Mg content is preferably 0.001% or more. The Mg content may be 0.005% or more, 0.010% or more, or 0.020% or more. On the other hand, if the Mg content exceeds 0.050%, coarse inclusions may form, which may reduce hydrogen embrittlement resistance. Therefore, the Mg content is set to 0.050% or less. The Mg content may also be 0.040% or less, 0.030% or less, or 0.020% or less.

[0039] Ca: 0 to 0.050% Ca is useful as a deoxidizing element and is also effective in controlling the morphology of sulfides. Therefore, Ca may be contained. To obtain the above effects, the Ca content is preferably 0.001% or more. The Ca content may be 0.002% or more, 0.004% or more, or 0.006% or more. On the other hand, if the Ca content exceeds 0.050%, coarse inclusions may be formed, which may reduce hydrogen embrittlement resistance. Therefore, the Ca content is set to 0.050% or less. The Ca content may be 0.040% or less, 0.030% or less, or 0.020% or less.

[0040] Y: 0 to 0.050% Y is an element that can control the morphology of sulfides when contained in a small amount, similar to Mg and Ca. Therefore, Y may be contained. To obtain the above effects, the Y content is preferably 0.001% or more. The Y content may be 0.002% or more, 0.004% or more, or 0.006% or more. On the other hand, if the Y content exceeds 0.050%, coarse Y oxides are formed, which may result in a decrease in hydrogen embrittlement resistance. Therefore, the Y content is set to 0.050% or less. The Y content may be 0.040% or less, 0.030% or less, or 0.020% or less.

[0041] Zr: 0 to 0.050% Zr is an element that can control the morphology of sulfides when contained in small amounts, similar to Mg, Ca, and Y. Therefore, Zr may be contained. To obtain the above effects, the Zr content is preferably 0.001% or more. The Zr content may be 0.002% or more, 0.004% or more, or 0.006% or more. On the other hand, if the Zr content exceeds 0.050%, coarse Zr oxides are formed, which may result in a decrease in hydrogen embrittlement resistance. Therefore, the Zr content is set to 0.050% or less. The Zr content may also be 0.040% or less, 0.030% or less, or 0.020% or less.

[0042] La: 0 to 0.050% Like Mg, Ca, Y, and Zr, La is an element that can control the morphology of sulfides even when contained in a small amount. Therefore, La may be contained. To obtain the above effects, the La content is preferably 0.001% or more. The La content may be 0.002% or more, 0.004% or more, or 0.006% or more. On the other hand, if the La content exceeds 0.050%, La oxides are formed, which may reduce hydrogen embrittlement resistance. Therefore, the La content is set to 0.050% or less. The La content may be 0.040% or less, 0.030% or less, or 0.020% or less.

[0043] Ce: 0 to 0.050% Ce, like La, is an element that can control the morphology of sulfides even when contained in a small amount. Therefore, Ce may be contained. To obtain the above effects, the Ce content is preferably 0.001% or more. The Ce content may be 0.002% or more, 0.004% or more, or 0.006% or more. On the other hand, if the Ce content exceeds 0.050%, Ce oxides are formed, which may reduce hydrogen embrittlement resistance. Therefore, the Ce content is set to 0.050% or less. The Ce content may also be 0.040% or less, 0.030% or less, or 0.020% or less.

[0044] Sn: 0 to 0.050% Sn is an element contained in steel when scrap is used as a raw material. If the Sn content is high, the hydrogen embrittlement resistance may be reduced due to embrittlement of grain boundaries. This adverse effect becomes particularly pronounced when the Sn content exceeds 0.050%. For this reason, the Sn content is set to 0.050% or less. The Sn content may also be 0.040% or less, 0.030% or less, or 0.020% or less. The lower the Sn content, the better, and even 0% is acceptable. However, if the Sn content is less than 0.001%, the refining cost increases. Therefore, the Sn content may be 0.001% or more. The Sn content may be 0.002% or more, 0.005% or more, or 0.010% or more.

[0045] Sb: 0 to 0.050% Like Sn, Sb is an element that is contained when scrap is used as a steel raw material. Sb is an element that strongly segregates at grain boundaries, causing embrittlement of the grain boundaries and a decrease in ductility. This adverse effect becomes particularly pronounced when the Sb content exceeds 0.050%. For this reason, the Sb content is set to 0.050% or less. The Sb content may also be 0.040% or less, 0.030% or less, or 0.020% or less. The lower the Sb content, the better, and even 0% is acceptable. However, if the Sb content is less than 0.001%, the refining cost increases. Therefore, the Sb content may be 0.001% or more. The Sb content may be 0.002% or more, 0.005% or more, or 0.008% or more.

[0046] As: 0 to 0.050% Like Sn and Sb, As is contained when scrap is used as a steel raw material. It is an element that strongly segregates at grain boundaries, causing embrittlement of the grain boundaries and reduced ductility. If the As content is high, hydrogen embrittlement resistance may be reduced. This adverse effect becomes particularly pronounced when the As content exceeds 0.050%. For this reason, the As content is set to 0.050% or less. The As content may also be 0.040% or less, 0.030% or less, or 0.020% or less. The lower the As content, the better, and 0% is acceptable. However, if the As content is less than 0.001%, the refining cost increases. Therefore, the As content may be 0.001% or more. The As content may be 0.002% or more, 0.003% or more, or 0.005% or more.

[0047] As described above, the chemical composition of the steel sheet according to this embodiment may contain the basic components with the balance being Fe and impurities, or may contain the basic components and further contain one or more optional components with the balance being Fe and impurities.

[0048] The chemical composition of the steel sheet according to this embodiment may be measured by a common method. For example, it may be measured using ICP-AES (Inductively Coupled Plasma-Atomic Emission Spectrometry) of chips in accordance with JIS G1201:2014. In this case, the chemical composition is the average content across the entire sheet thickness. C and S, which cannot be measured by ICP-AES, may be measured using a combustion-infrared absorption method, N may be measured using an inert gas fusion-thermal conductivity method, and O may be measured using an inert gas fusion-non-dispersive infrared absorption method. The Mn content determined here is referred to as [Mnave], which will be described later. When the steel sheet has a coating layer on its surface, the coating layer can be removed by mechanical grinding or the like before analyzing the chemical composition. When the coating layer is a plating layer, the plating layer can be removed by dissolving it in an acid solution to which an inhibitor that suppresses corrosion of the steel sheet has been added.

[0049] <Microstructure (metal structure)> Next, the microstructure of the steel sheet according to this embodiment will be described. In this embodiment, the microstructure is the microstructure located in the range of 1 / 8 to 3 / 8 of the sheet thickness from the surface of the steel sheet in the sheet thickness direction (t / 4 part). The microstructure of the t / 4 part is specified because it is a representative microstructure of the steel sheet and has a high correlation with the properties of the steel sheet. The percentage (%) of each phase below is an area percentage unless otherwise specified.

[0050] Ferrite: 5.0% or less Ferrite affects the deformability of steels whose main structure is martensite and tempered martensite. As the area fraction of ferrite increases, local deformability and hydrogen embrittlement resistance decrease. In particular, if the area fraction of ferrite exceeds 5.0%, fracture during elastic deformation under stress load may decrease hydrogen embrittlement resistance. Therefore, the area fraction of ferrite is set to 5.0% or less. The area fraction of ferrite may also be 4.0% or less, 3.0% or less, or 2.0% or less. The area ratio of ferrite may be 0%, but to make it less than 1.0%, advanced control is required in the manufacturing process, which leads to a decrease in yield. Therefore, the area ratio of ferrite may be set to 1.0% or more.

[0051] Martensite and tempered martensite: total of more than 90.0% The total area ratio of martensite (so-called fresh martensite) and tempered martensite affects the strength of the steel, with a higher area ratio resulting in increased tensile strength. If the total area ratio of martensite and tempered martensite is 90.0% or less, the target tensile strength cannot be achieved, and this may also result in fracture during elastic deformation under stress load and a decrease in hydrogen embrittlement resistance. For this reason, the total area ratio of martensite and tempered martensite is set to more than 90.0%. The total area ratio of martensite and tempered martensite may be 95.0% or more, 97.0% or more, 99.0% or more, or 100.0%.

[0052] Remainder: Consists of one or more of bainite, pearlite, and retained austenite The area ratio of the structure other than the above structure (remaining structure) may be 0%, but if a remaining structure exists, the remaining structure is composed of one or more of bainite, pearlite, and retained austenite. If the area ratio of the remaining structure exceeds 8.0%, fracture may occur during elastic deformation under stress load, and hydrogen embrittlement resistance may decrease. Therefore, the area ratio of the remaining structure is preferably 8.0% or less, and more preferably 7.0% or less. Among these, pearlite and retained austenite are particularly structures that deteriorate the local ductility of steel, so the smaller the amount, the better. On the other hand, achieving a 0% area ratio of the remaining structure requires advanced control in manufacturing, which may result in a decrease in yield, so the area ratio of the remaining structure may be 1.0% or more.

[0053] The area ratio of each phase in the microstructure of the steel sheet according to this embodiment can be determined by the following method.

[0054] (Method for evaluating the area ratio of ferrite) The ferrite area fraction was determined by observing the t / 4 region (a region from 1 / 8 to 3 / 8 of the plate thickness, centered at 1 / 4 of the plate thickness from the surface in the plate thickness direction) using electron channeling contrast images taken with a field emission scanning electron microscope (FE-SEM). Electron channeling contrast images are a method for detecting crystal orientation differences within grains as differences in image contrast. In these images, polygonal ferrite is defined as the region that appears with uniform contrast in a structure determined to be ferrite, rather than pearlite, bainite, martensite, or retained austenite. Eight 35 μm × 25 μm electron channeling contrast images were analyzed using image analysis to calculate the polygonal ferrite area fraction for each field, and the average value was used to determine the ferrite area fraction.

[0055] (Method for evaluating the total area ratio of martensite and tempered martensite) The total area fraction of martensite and tempered martensite is also determined from images captured using the electron channeling contrast technique described above. Because these structures are less susceptible to etching than ferrite, they appear as convexities on the microstructure observation surface. Tempered martensite is a collection of lath-shaped crystal grains containing iron-based carbides with a long diameter of 20 nm or more. These carbides belong to multiple variants, i.e., multiple iron-based carbide groups elongated in different directions. Retained austenite also appears as convexities on the microstructure observation surface. Therefore, by subtracting the area fraction of the convexities determined using the above procedure from the area fraction of retained austenite measured using the procedure described below, the total area fraction of martensite and tempered martensite can be accurately measured.

[0056] (Method for evaluating the total area ratio of bainite, pearlite and retained austenite) The area fraction of retained austenite can be calculated by X-ray measurement. Specifically, the specimen is mechanically and chemically polished from the surface to a position one-quarter of the thickness in the thickness direction. The polished specimen is then subjected to MoKα X-ray analysis. The MoKα X-rays are then used as characteristic X-rays to measure the integrated intensity ratio of the (200), (211) diffraction peaks of the bcc phase and the (200), (220), and (311) diffraction peaks of the fcc phase. The fraction of retained austenite is calculated from the ratio, and this is the area fraction of retained austenite. The area ratio of pearlite is determined from images taken using the electron channeling contrast method described above. Pearlite is a structure in which plate-like carbides and ferrite are arranged. Bainite is a collection of lath-shaped crystal grains that either does not contain iron-based carbides with a major axis of 20 nm or more, or contains iron-based carbides with a major axis of 20 nm or more, but the carbides belong to a single variant, i.e., a group of iron-based carbides elongated in the same direction. Here, the group of iron-based carbides elongated in the same direction refers to iron-based carbides whose elongation directions differ by no more than 5°.

[0057] (The area where the Mn content is 1.1 × [Mnave] or more has a number density of 5.0 × 10 -4 pieces / μm 2or more, and the distance to the nearest region where the Mn content is 1.1 × [Mnave] or more is 10.0 μm or less on average) Since Mn has an attractive interaction with H, controlling its dispersion state can suppress the accumulation of H at the prior austenite grain boundaries, resulting in the production of steel sheets with excellent hydrogen embrittlement resistance. In the steel sheet according to this embodiment, a plurality of regions having a Mn content of 1.1 × [Mnave] or more are dispersed in a cross section in the sheet thickness direction as regions having a stronger attractive interaction with H than the prior austenite grain boundary, where [Mnave] is the average Mn content throughout the sheet thickness direction. On the other hand, if the distance between these multiple regions, i.e., the distance between a region where the Mn content is 1.1 × [Mnave] or more and the nearest other region where the Mn content is 1.1 × [Mnave] or more, exceeds 10.0 μm on average, the effect of suppressing H segregation to the prior austenite grain boundaries cannot be sufficiently obtained, and hydrogen embrittlement resistance decreases. Therefore, a plurality of regions having a Mn content of 1.1×[Mnave] or more are present (dispersed) so that the distance between the nearest regions having a Mn content of 1.1×[Mnave] or more is 10.0 μm or less on average. Furthermore, even if the regions with a Mn content of 1.1 × [Mnave] or more are dispersed as described above, if the number density is low, sufficient effect cannot be obtained. -4 pieces / μm 2 That's all.

[0058] The identification of regions where the Mn content is 1.1 × [Mnave] or more, the number density thereof, and the average distance between adjacent regions are determined by the following method. The Mn dispersion state was measured using an electron probe microanalyzer (EPMA). Specifically, a sample was taken so that the cross section parallel to the rolling direction of the steel sheet served as the measurement surface. A 50 μm × 50 μm area was defined as one field of view in the t / 4 portion of the observation surface (a range of 1 / 8 to 3 / 8 of the sheet thickness, centered at 1 / 4 of the sheet thickness in the thickness direction from the surface). The Mn element concentration was then mapped at a measurement interval of 0.1 μm. This was performed for 10 fields of view, and the elemental map data for the 10 fields of view was collected as numerical data. A binary image was then created, color-coded to indicate areas with a Mn concentration of 1.1 × [Mnave] or greater and areas with a Mn concentration less than 1.1 × [Mnave]. In this binary image, areas where multiple pixels are connected and the Mn concentration is 1.1 × [Mnave] or greater were considered to be a single area, and the number of areas with a Mn concentration of 1.1 × [Mnave] or greater was calculated. The number of particles obtained is divided by the area of one field of view to determine the number density of the region where the number is 1.1 × [Mnave] or more. The area of one field of view is divided by the number obtained as described above, and the result is raised to the power of 1 / 2 to determine the average interval between the regions where the number is 1.1 × [Mnave] or more.

[0059] (mechanical properties) The steel sheet according to this embodiment has a tensile strength (TS) of 1500 MPa or more, which is a strength that contributes to reducing the weight of an automobile body. There is no need to set an upper limit, but since an increase in tensile strength may result in a decrease in formability, the tensile strength may be set to 2500 MPa or less, or 2000 MPa or less.

[0060] (plate thickness) The thickness of the steel sheet according to this embodiment is not limited, but is preferably 1.0 mm or more and 2.2 mm or less. The thickness is more preferably 1.05 mm or more or 1.1 mm or more. The thickness is more preferably 2.1 mm or less or 2.0 mm or less.

[0061] (Coating layer) The steel sheet according to this embodiment may have a coating layer containing zinc, aluminum, magnesium or an alloy thereof on one or both surfaces, and the coating layer may consist of zinc, aluminum, magnesium or an alloy thereof and impurities. By providing a coating layer on the surface, corrosion resistance is improved. In some cases, automotive steel sheets cannot be made thinner than a certain thickness even if they are made stronger due to concerns about holes caused by corrosion. One of the goals of increasing the strength of steel sheets is to reduce weight by making them thinner, so even if high-strength steel sheets are developed, if their corrosion resistance is low, the areas in which they can be used will be limited. One method that can be considered to solve these issues is to form a coating layer on both the front and back surfaces to improve corrosion resistance. Even if the coating layer is formed, the hydrogen embrittlement resistance of the steel plate according to this embodiment is not impaired. The coating layer is, for example, a hot-dip galvanized layer, a galvannealed layer, an electrolytic galvanized layer, an aluminum plated layer, a Zn-Al alloy plated layer, an Al-Mg alloy plated layer, or a Zn-Al-Mg alloy plated layer. When a coating layer is present on the surface, the surface that serves as the reference for the t / 4 portion is the surface of the base steel excluding the coating layer.

[0062] <Manufacturing method> A manufacturing method capable of manufacturing the steel sheet according to the present embodiment (a manufacturing method of the steel sheet according to the present embodiment) will be described. The method for manufacturing a steel sheet according to this embodiment is as follows: (I) a heating step of heating a steel billet having a predetermined chemical composition; (II) a hot rolling step of hot rolling the heated steel slab to obtain a hot-rolled steel sheet; (III) a post-hot rolling cooling step in which cooling of the hot-rolled steel sheet is started less than 1.0 second after the completion of the hot rolling step, and cooled to a coiling temperature of 400°C or more and less than 550°C at an average cooling rate of 20°C / second or more and 50°C / second or less; (IV) a coiling step of coiling the hot-rolled steel sheet after the post-hot-rolling cooling step at the coiling temperature; (V) a cold rolling step of subjecting the hot-rolled steel sheet after the coiling step to pickling and cold rolling to obtain a cold-rolled steel sheet; (VI) an annealing step of heating the cold-rolled steel sheet after the cold rolling step to an annealing temperature at an average heating rate of 15°C / sec or more and 100°C / sec or less from room temperature to 700°C, and at an average heating rate of 5°C / sec or more and less than 15°C / sec from 700°C to an annealing temperature of 830°C or more and less than 900°C, and holding the annealing temperature for 25 to 100 seconds to anneal the steel sheet; (VII) a post-annealing cooling process in which the cold-rolled steel sheet after the annealing process is cooled to 25 to 300 ° C. at an average cooling rate of 4 to 100 ° C. / second; Equipped with. Preferred conditions for each step will be explained below.

[0063] [Heating process] In the heating step, a steel piece such as a slab having a predetermined chemical composition (the chemical composition is substantially unchanged during the manufacturing stage, and therefore has the same chemical composition as the steel sheet according to this embodiment) is heated prior to the hot rolling step. The heating temperature is not limited as long as the rolling temperature in the next step can be ensured, for example, 1000 to 1300°C. From the viewpoint of productivity, the steel billets used are preferably cast by continuous casting, but may also be produced by ingot casting or thin slab casting. If the steel slab obtained by continuous casting can be subjected to the hot rolling step while still at a sufficiently high temperature, the heating step may be omitted.

[0064] [Hot rolling process] In the hot rolling step, the heated steel billet is hot rolled to obtain a hot-rolled steel sheet. The hot rolling process includes rough rolling and finish rolling, and the finish rolling involves multiple passes of reduction, with at least four passes being large reduction passes with a reduction rate of 20% or more, and the inter-pass time between the large reduction passes being 5.0 seconds or less. The rolling start temperature is 950 to 1100°C, and the rolling end temperature is 800 to 950°C.

[0065] (In finish rolling, large reduction passes with a reduction rate of 20% or more: 4 passes or more) (Interpass time for large reduction passes: within 5.0 seconds) By controlling the reduction ratio, number of rolling passes, and inter-pass time in finish rolling, it is possible to control the morphology of austenite grains to be equiaxed and fine. When the austenite grains become equiaxed and fine, a uniform and fine pearlite structure is formed in the subsequent cooling process after hot rolling. Since the pearlite structure contains cementite, it is possible to obtain a structure in which cementite is uniformly and finely precipitated (dispersed precipitation). If there are fewer than four passes (large reduction passes) with a reduction ratio of 20% or more, unrecrystallized austenite remains, and sufficient effect cannot be obtained. Therefore, for four or more passes, the reduction ratio is set to 20% or more (four or more passes are performed with a reduction ratio of 20% or more). Preferably, for five or more passes, the reduction ratio is set to 20% or more. On the other hand, there is no particular upper limit on the number of passes with a reduction ratio of 20% or more, but for more than 10 passes, a large number of rolling stands must be installed, which may result in an increase in the size of the equipment and manufacturing costs. Therefore, the number of passes with a rolling reduction of 20% or more (number of passes) may be 10 passes or less, 9 passes or less, or 7 passes or less. Furthermore, the interpass time in finish rolling has a significant effect on the recrystallization and grain growth of austenite grains after rolling. Even when four or more large reduction passes are used, if the interpass time between each large reduction pass exceeds 5.0 seconds, grain growth is likely to occur, causing the austenite grains to coarsen. Preferably, the interpass time is 3.0 seconds or less, or 1.0 second or less. On the other hand, although there is no need to set a lower limit for the interpass time, if the interpass time between heavy reduction passes is less than 0.2 seconds, the recrystallization of austenite may not be completed, the proportion of unrecrystallized austenite may increase, and sufficient effects may not be obtained. For this reason, it is preferable that the interpass time between heavy reduction passes is 0.2 seconds or more. The interpass time may be 0.3 seconds or more, or 0.5 seconds or more.

[0066] (Rolling start temperature: 950~1100℃) (Rolling end temperature: 800-950℃) The rolling start temperature is an important factor in controlling austenite recrystallization. If the rolling start temperature is less than 950°C, the temperature drops during rolling, leaving unrecrystallized austenite. This causes ferrite grains to form along the elongated unrecrystallized austenite grain boundaries, and the untransformed austenite within the grains becomes pearlite. In this case, the size of the ferrite grains increases, and when Mn is concentrated in the pearlite structure, the average spacing of the Mn-enriched areas exceeds 10.0 μm. On the other hand, if the rolling start temperature exceeds 1100°C, the temperature during rolling becomes so high that alloying elements that inhibit ferrite transformation tend to concentrate at the austenite grain boundaries. In this case, the ferrite transformation is delayed during the cooling process after finish rolling, the proportion of pearlite increases, and Mn concentration in cementite in the pearlite structure becomes insufficient. In this case, it is not possible to obtain a sufficient region where the Mn content is 1.1 × [Mnave] or more. Furthermore, if the rolling end temperature is less than 800°C, unrecrystallized austenite remains, and as a result, the average interval of the domains having a value of 1.1×[Mnave] becomes large. Furthermore, if the rolling end temperature exceeds 950°C, the ferrite transformation is excessively suppressed, the intervals between cementite grains where Mn is concentrated become larger, and the number density decreases. As a result, the average intervals and number density of the regions where the Mn content is 1.1 × [Mnave] or more do not fall within the preferred ranges. Therefore, the rolling start temperature is set to 950 to 1100°C, and the rolling end temperature is set to 800 to 950°C.

[0067] [Cooling process after hot rolling] [Winding process] In the post-hot rolling cooling step, cooling of the hot-rolled steel sheet obtained in the hot rolling step is started less than 1.0 second after the completion of the hot rolling step, and is cooled at an average cooling rate of 20°C / second to 50°C / second to a coiling temperature of 400°C to less than 550°C. In the coiling step, the hot-rolled steel sheet after the post-hot rolling cooling step is coiled at the above-mentioned coiling temperature. By performing these steps under predetermined conditions, the pearlite structure is finely dispersed during cooling, and accordingly, the cementite in the pearlite structure is also finely dispersed. Furthermore, Mn is concentrated in the cementite after coiling. If the time from the end of hot rolling to the start of cooling is 1.0 second or longer, the average cooling rate is less than 20°C / second, or the cooling stop temperature (coiling temperature) is 550°C or higher, excessive grain growth of ferrite grains occurs, increasing the dispersion intervals between pearlite structures, which undesirably widens the average intervals of cementite where Mn is concentrated after coiling. On the other hand, if the average cooling rate exceeds 50°C / sec, hardening phases such as bainite and martensite are likely to form, and the proportion of pearlite structure decreases, making it impossible to obtain a sufficient amount of cementite with concentrated Mn. Furthermore, if the cooling stop temperature is less than 400°C, Mn may not be sufficiently concentrated in the cementite in the pearlite structure.

[0068] [Cold rolling process] In the cold rolling step, the hot rolled steel sheet after the coiling step is uncoiled, pickled and cold rolled to obtain a cold rolled steel sheet. Pickling removes oxide scale from the surface of the hot-rolled steel sheet, improving the chemical conversion treatability and plating properties of the cold-rolled steel sheet. Pickling may be performed under known conditions, and may be performed once or multiple times. The reduction ratio of cold rolling is not particularly limited, and is, for example, 20 to 80%.

[0069] [Annealing process] [Cooling process after annealing] In the annealing step, the cold-rolled steel sheet after the cold rolling step is heated to an annealing temperature at an average rate of 15 to 100°C / sec from room temperature (e.g., 25°C) to 700°C, and at an average rate of 5°C / sec to less than 15°C / sec from 700°C to an annealing temperature of 830°C or higher and 900°C, and is then annealed by holding the annealing temperature for 25 to 100 seconds. The average rate of temperature rise to 700°C is preferably 17°C / sec or higher, more preferably 20°C / sec or higher. The average rate of temperature rise from 700°C to an annealing temperature of 830°C or higher and lower than 900°C is preferably 6°C / sec or higher and 14°C / sec or lower, more preferably 7°C / sec or higher and 13°C / sec or lower. Thereafter, in the post-annealing cooling step, the cold-rolled steel sheet after the annealing step is cooled to 25 to 300°C at an average cooling rate of 4 to 100°C / sec. In these processes, the steel is heated to the austenite single-phase region while maintaining the dispersion of concentrated Mn in the cementite, dissolving the cementite, and then cooled to below 300°C, resulting in a martensite-based structure with dispersed Mn-enriched areas.

[0070] During the temperature rise process in the annealing step, if the average temperature rise rate from room temperature (e.g., 25°C) to 700°C is less than 15°C / s, cementite will coarsen. Since the size of cementite is the same as that of the Mn-enriched region, the average spacing of regions that are 1.1 × [Mnave] or greater after annealing will increase and the number density may decrease. On the other hand, an average temperature rise rate of more than 100°C / s requires special equipment, which increases practical costs. Furthermore, in the temperature range from 700°C to the annealing temperature (830°C or higher, but less than 900°C), recrystallization is promoted and the density of grain boundaries and dislocations is reduced, thereby suppressing Mn diffusion. If the average heating rate in this temperature range is less than 5°C / sec, Mn may diffuse, resulting in the dispersion of Mn-enriched regions being eliminated. On the other hand, if the average heating rate is 15°C / sec or higher, recrystallization does not occur, and Mn may diffuse through grain boundaries and dislocations, resulting in the dispersion of Mn-enriched regions being eliminated. If the annealing temperature is less than 830°C, the austenite reverse transformation will not be completed, the volume fraction of the ferrite structure will increase, and the strength may not be achieved. This is because the Mn-enriched cementite has high thermal stability and is likely to remain without dissolving, which in turn causes the ferrite structure to remain. On the other hand, if the annealing temperature is 900°C or higher, Mn will diffuse, and the well-dispersed Mn-enriched parts will be lost. If the annealing time (holding time) is less than 25 seconds, austenitization may be insufficient and the desired structure may not be obtained, whereas if the annealing time exceeds 100 seconds, Mn may diffuse, resulting in the dispersion of Mn-enriched parts being eliminated.

[0071] In the post-annealing cooling process, if the average cooling rate to the cooling stop temperature is less than 4°C / s, structures such as ferrite and bainite will be formed, making it impossible to obtain the desired structure. On the other hand, if the average cooling rate exceeds 100°C / s, special equipment will be required, which will increase the practical cost. If the cooling stop temperature exceeds 300°C, a bainite structure will form, making it impossible to obtain the desired structure. On the other hand, if the cooling stop temperature is set to less than 25°C, a special refrigerant or the like will need to be used, which will cause problems in terms of cost and productivity.

[0072] [Tempering process] In the method for manufacturing a steel sheet according to this embodiment, the cold-rolled steel sheet after the annealing step may further be subjected to a tempering step in which the temperature is increased to 50°C or more and less than 500°C and held for 5 seconds or more and less than 1000 seconds. By performing tempering under the above conditions, martensite is converted into tempered martensite, thereby improving formability. If the tempering temperature (holding temperature) is less than 50°C or the holding time is less than 5 seconds, the above effects cannot be obtained. On the other hand, if the tempering temperature is 500°C or higher, Mn may diffuse, causing the dispersion of Mn-enriched areas to disappear. In addition, a decrease in dislocation density in the tempered martensite may occur, resulting in a decrease in strength and a decrease in tensile strength. Furthermore, if the holding time is 1000 seconds or longer, not only will the strength decrease, but productivity will also decrease. Tempering may be performed in continuous annealing equipment, or it may be performed offline after continuous annealing in a separate equipment.

[0073] In the method for producing a steel sheet according to this embodiment, a coating layer containing zinc, aluminum, magnesium or an alloy thereof may be formed on the surface of the steel sheet in any of the steps from the post-annealing cooling step to the tempering step. The coating layer is preferably a coating layer containing zinc, aluminum, magnesium or an alloy thereof. The coating layer is, for example, a plating layer. The coating method is not limited, but for example, when forming a coating layer mainly composed of zinc by hot dip plating, an example of conditions for forming the plating layer is to adjust (by heating or cooling) the steel sheet temperature to (plating bath temperature -40)°C to (plating bath temperature +50)°C, and then immerse the cold rolled steel sheet in a plating bath at 450 to 490°C. The reason this condition is preferable is that if the temperature of the steel sheet during immersion in the coating bath is below the hot-dip galvanizing bath temperature -40°C, heat loss during immersion is large, which can cause some of the molten zinc to solidify and deteriorate the coating appearance, and if the temperature exceeds the hot-dip galvanizing bath temperature +50°C, operational problems can occur due to the rise in the coating bath temperature. When forming a zinc-based coating layer, the coating bath preferably has an effective Al content (the total Al content minus the total Fe content in the coating bath) of 0.050 to 0.250 mass%, optionally containing Mg, with the remainder being Zn and impurities. If the effective Al content in the coating bath is less than 0.050 mass%, excessive penetration of Fe into the coating layer may occur, resulting in reduced coating adhesion. On the other hand, if the effective Al content in the coating bath exceeds 0.250 mass%, Al-based oxides that inhibit the migration of Fe and Zn atoms may form at the boundary between the steel sheet and the coating layer, potentially reducing coating adhesion. [Example]

[0074] [Example 1] Steels having the chemical compositions shown in Tables 1-1 and 1-2 were melted and cast into billets. The billets were placed in a furnace heated to 1200°C and subjected to a homogenization treatment by holding them there for 60 minutes. They were then removed from the furnace and hot-rolled to obtain steel plates with a thickness of 2.8 mm. In the hot-rolling process, a rolling mill with seven stands was used to continuously perform seven passes of finish rolling (so that the inter-pass time was constant), with four passes resulting in a rolling reduction of more than 20%. The inter-pass time between each pass resulting in a 20% or greater reduction in the finish rolling and the pass immediately preceding that pass was 0.6 seconds. The start temperature of the finish rolling was 1070°C, the end temperature was 850°C, and 0.8 seconds after the end of the finish rolling, the billets were water-cooled to 530°C at an average cooling rate of 29.0°C / s, and then coiled. However, for AH-0, the steel plate became embrittled and cracked during hot rolling, so no further processes were carried out. Subsequently, the oxide scale of this hot-rolled steel sheet was removed by pickling, and the sheet was cold-rolled at a reduction ratio of 50.0% to a finished sheet thickness of 1.4 mm. Furthermore, this cold-rolled steel sheet was heated from room temperature to 700°C at an average heating rate of 35.0°C / s, and then heated from 700°C to 860°C at an average heating rate of 10°C / s. After being held at 860°C for 80 seconds, it was cooled to 190°C at an average cooling rate of 38.0°C / s. Subsequently, the specimen was reheated to 230°C and held there for 180 seconds for tempering. No plating treatment was performed.

[0075] The microstructure of the obtained cold-rolled steel sheet was observed as described above, and the area ratio of each phase in the quarter thickness part was determined. In addition, the number density of regions in the quarter thickness part where the Mn content was 1.1 × [Mnave] or more and the average distance between the nearest other regions were measured. The results are shown in Table 2. Furthermore, the chemical composition of the samples collected from the manufactured steel plates was analyzed and found to be the same as the chemical composition of the steel shown in Tables 1-1 and 1-2.

[0076] The obtained cold-rolled steel sheets were evaluated for tensile properties and hydrogen embrittlement resistance (hydrogen embrittlement resistance) in the following manner.

[0077] (Method for evaluating tensile properties) The tensile test was conducted in accordance with JIS Z 2241 (2011), using JIS No. 5 test pieces taken in a direction in which the longitudinal direction of the test piece was parallel to the rolling direction of the steel strip, and the tensile strength (TS) and total elongation (El) were measured.

[0078] (Method for evaluating hydrogen embrittlement resistance) The hydrogen embrittlement resistance of steel sheets manufactured using the steel sheet manufacturing method according to an embodiment of the present invention was evaluated by the following method. Specifically, the steel sheets were sheared at a clearance of 10%, and then subjected to a U-bend test at 10R. A strain gauge was attached to the center of the obtained test piece, and both ends of the test piece were fastened with bolts to apply stress. The applied stress was calculated from the strain monitored by the strain gauge. A load stress corresponding to 80% of the tensile strength (TS) was applied (for example, in the case of A-0 in Table 2, the applied stress = 2213 MPa × 0.8 = 1770 MPa). This is because it is believed that the residual stress introduced during forming corresponds to the tensile strength of the steel sheet. The obtained U-bend test pieces were immersed in an HCl aqueous solution with a pH of 2 at a liquid temperature of 25°C and held for 48 hours, after which the presence or absence of cracks was examined. The lower the pH of the HCl aqueous solution and the longer the immersion time, the greater the amount of hydrogen that penetrates into the steel sheet, resulting in a severe hydrogen embrittlement environment. After immersion, if a crack longer than 1.00 mm was observed in the U-bend test piece, it was evaluated as NG, and if no crack longer than 1.00 mm was observed, it was evaluated as OK.

[0079] Steel sheets with a tensile strength of 1500 MPa or more and with hydrogen embrittlement resistance evaluated as OK were evaluated as having high strength and excellent hydrogen embrittlement resistance.

[0080] [Table 1-1]

[0081] [Table 1-2]

[0082] [Table 2]

[0083] As can be seen from Tables 1-1 and 1-2, Nos. A-0 to O-0 had chemical compositions, area ratios of microstructures, number densities of regions where the Mn content was 1.1 × [Mnave] or more, and average intervals between the nearest other regions all within the ranges of the present invention, and were excellent in tensile strength and hydrogen embrittlement resistance.

[0084] In contrast, Nos. P-0 to AN-0 had chemical compositions outside the range of the present invention, and were therefore inferior in one or more of tensile strength and hydrogen embrittlement resistance.

[0085] P-0 had a tensile strength of less than 1500 MPa due to its low C content. Q-0 had a high C content, which resulted in a decrease in hydrogen embrittlement resistance. R-0 had a high Si content, which resulted in a decrease in hydrogen embrittlement resistance. The tensile strength of S-0 was less than 1500 MPa due to the low Mn content. T-0 had a high Mn content, which resulted in poor hydrogen embrittlement resistance. The high P content of U-0 resulted in a decrease in hydrogen embrittlement resistance due to grain boundary embrittlement. V-0 had a high S content, which resulted in a decrease in hydrogen embrittlement resistance. W-0 had a high Al content, which resulted in the formation of coarse Al oxides, which reduced the hydrogen embrittlement resistance. X-0 had a high N content, which resulted in the formation of coarse nitrides, reducing its hydrogen embrittlement resistance.

[0086] Since Y-0 had a high O content, oxides were formed, which reduced the hydrogen embrittlement resistance. Z-0 had a high Co content, which resulted in the precipitation of coarse Co carbides, which reduced hydrogen embrittlement resistance. AA-0 had a high Ni content, which resulted in a decrease in hydrogen embrittlement resistance. AB-0 had a high Mo content, which resulted in the crystallization of coarse Mo carbides, which reduced hydrogen embrittlement resistance. AC-0 had a high Cr content, which resulted in the precipitation of coarse Cr carbides, which reduced its hydrogen embrittlement resistance. AD-0 had a high Ti content, which resulted in increased precipitation of carbonitrides and reduced hydrogen embrittlement resistance. AE-0 had a high B content, which resulted in the formation of coarse B oxides in the steel, which reduced its hydrogen embrittlement resistance. AF-0 had a high Nb content, which resulted in the formation of coarse Nb carbides, which reduced the hydrogen embrittlement resistance. AG-0 had a high V content, which resulted in increased precipitation of carbonitrides and reduced hydrogen embrittlement resistance.

[0087] Because AH-0 had a high Cu content, the steel plate became embrittled and cracked during hot rolling, and no further evaluation was carried out. Since AI-0 had a high W content, coarse W precipitates were formed, which reduced the hydrogen embrittlement resistance. AJ-0 and AK-0 had high Mg and Ca contents, which resulted in the formation of coarse inclusions and reduced hydrogen embrittlement resistance. AL-0 had a high Zr content, which resulted in the formation of coarse Zr oxides, which reduced the hydrogen embrittlement resistance. The high Sn and Sb contents of AM-0 and AN-0, respectively, led to a decrease in hydrogen embrittlement resistance due to grain boundary segregation.

[0088] [Example 2] Furthermore, to investigate the influence of manufacturing conditions, hot-rolled steel sheets with a thickness of 2.3 mm were produced using the same equipment as in Example 1 under the manufacturing conditions listed in Table 3-1 for steel types A to O, which were found to have excellent properties in Table 2. These hot-rolled steel sheets were then cold-rolled at a rolling reduction of 55% to obtain cold-rolled steel sheets, which were then annealed, cooled after annealing, and tempered as needed under the conditions listed in Tables 3-2 and 3-3. Some of the cold-rolled steel sheets were also plated to form a zinc-coated layer on their surfaces. The designations GI and GA for the plating types in Table 3-3 indicate the zinc-coating method. GI is a steel sheet obtained by immersing the steel sheet in a hot-dip galvanizing bath at 465°C to form a zinc-coated layer on its surface, and GA is a steel sheet obtained by immersing the steel sheet in a hot-dip galvanizing bath at 465°C and then heating the steel sheet to 490°C to form an iron-zinc alloy layer on its surface. In Table 3-3, examples in which tempering is marked with "-" are examples in which tempering was not performed.

[0089] The microstructure of the obtained cold-rolled steel sheet was observed to determine the area ratio of each phase in the microstructure of the quarter thickness portion in the same manner as in Example 1. Furthermore, the number density of regions in the quarter thickness portion where the Mn content was 1.1 × [Mnave] or more and the average distance between the regions and the nearest other regions where the Mn content was 1.1 × [Mnave] or more were measured. The tensile properties of the obtained cold-rolled steel sheets were evaluated in the same manner as in Example 1. The hydrogen embrittlement resistance (hydrogen embrittlement resistance characteristics) was evaluated by the following method.

[0090] (Method for evaluating hydrogen embrittlement resistance) After shearing the steel plate with a clearance of 10%, a U-bend test was performed at 10R. A strain gauge was attached to the center of the obtained test piece, and stress was applied by fastening both ends of the test piece with bolts. The applied stress was calculated from the strain monitored by the strain gauge. The applied stress was a stress corresponding to 80% of the tensile strength (TS) (for example, in the case of A-1 in Table 4, the applied stress = 2101 MPa x 0.8 = 1681 MPa). This is because the residual stress introduced during forming is thought to correspond to the tensile strength of the steel plate. The obtained U-bend test specimens were immersed in an HCl aqueous solution with a pH of 2 at a liquid temperature of 25°C and held for 96 hours, after which the presence or absence of cracks was examined. The lower the pH of the HCl aqueous solution and the longer the immersion time, the greater the amount of hydrogen that penetrates into the steel sheet, resulting in a more severe hydrogen embrittlement environment. After immersion, the total length of the cracks in the U-bend test specimens was measured (if multiple cracks were found, the total length was calculated as the sum of the individual measurements). The smaller the total crack length, the better the hydrogen embrittlement resistance. However, when cracks longer than 1.00 mm were observed, the specimen was rated as NG, and when no cracks or minor cracks with a length of 1.00 mm or less were observed, the specimen was rated as OK. An OK rating was used to indicate passing, and an NG rating was used to indicate failing. When the crack length was 0.50 mm or less, the specimen was judged to have particularly excellent hydrogen embrittlement resistance. The results obtained are shown in Table 4.

[0091] [Table 3-1]

[0092] [Table 3-2]

[0093] [Table 3-3]

[0094] [Table 4]

[0095] As can be seen from Tables 3-1 to 3-3 and Table 4, in all the examples according to the present invention, by appropriately controlling the conditions of hot rolling, coiling and annealing, and cooling after annealing, in particular, steel sheets with high strength and excellent hydrogen embrittlement resistance could be obtained.

[0096] On the other hand, in A-2, the hot rolling start temperature was low, so the temperature dropped during rolling and a large amount of unrecrystallized austenite remained. As a result, the average spacing of the 1.1 × [Mnave] region increased, and hydrogen embrittlement resistance deteriorated. Because the hot rolling start temperature of B-2 was high, alloying elements segregated at the austenite grain boundaries during rolling, delaying the ferrite transformation. As a result, the number density of the 1.1 × [Mnave] region decreased, and hydrogen embrittlement resistance deteriorated. In C-2, the hot rolling end temperature was low, so unrecrystallized austenite was formed, resulting in an increase in the average spacing of the 1.1 × [Mnave] domains and a deterioration in hydrogen embrittlement resistance. In D-2, the hot rolling end temperature was high, so the ferrite transformation was excessively suppressed, resulting in a decrease in the number density of the 1.1 × [Mnave] region and a deterioration in hydrogen embrittlement resistance. In E-2, the interpass time of hot rolling was long, so a large amount of unrecrystallized austenite remained. As a result, the average spacing of the 1.1 × [Mnave] region became larger, and hydrogen embrittlement resistance deteriorated. In F-2, the number of passes at a reduction rate of 20% or more was small, resulting in a larger average interval between regions with 1.1 × [Mnave] and a deterioration in hydrogen embrittlement resistance. In H-2, the time from finish rolling to the start of cooling was long, which caused excessive ferrite transformation, increasing the average spacing of the 1.1 × [Mnave] region and deteriorating hydrogen embrittlement resistance. In I-2, the cooling rate after hot rolling was slow, which caused excessive ferrite transformation, increasing the average spacing of the 1.1 × [Mnave] region, and deteriorating hydrogen embrittlement resistance. In J-2, the cooling rate after the finish rolling was high, resulting in the formation of a large amount of bainite and martensite structures. As a result, the number density of the 1.1 × [Mnave] region decreased, and the hydrogen embrittlement resistance deteriorated. In K-2, the coiling temperature was low, so bainite and martensite structures were formed, resulting in a decrease in the number density of regions with 1.1 × [Mnave] and a deterioration in hydrogen embrittlement resistance. In L-2, the coiling temperature was high, which increased the proportion of ferrite structure. As a result, the average spacing of the 1.1 × [Mnave] region increased, and hydrogen embrittlement resistance deteriorated.

[0097] In M-2 and N-3, the average heating rate up to 700°C was slow, causing cementite to coarsen during the heating process, increasing the average spacing of the 1.1 × [Mnave] domains. In N-3, the number density of the 1.1 × [Mnave] domains also decreased. As a result, hydrogen embrittlement resistance deteriorated. In O-2 and N-4, the rate of temperature rise to the maximum heating temperature (annealing temperature) during the annealing process was slow, which caused Mn diffusion, reducing the number density of the 1.1 × [Mnave] region and deteriorating hydrogen embrittlement resistance. In A-3 and N-5, the heating rate to the maximum heating temperature during the annealing process was too fast, which prevented complete recrystallization. The Mn-enriched areas were eliminated by dislocations and grain boundaries, resulting in a decrease in the number density of the regions satisfying 1.1 × [Mnave]. Furthermore, in N-5, the average spacing of the regions satisfying 1.1 × [Mnave] also increased. As a result, hydrogen embrittlement resistance decreased. In B-3, the maximum heating temperature during the annealing process was low, so the proportion of ferrite exceeded 5%, resulting in poor hydrogen embrittlement resistance. In addition, because the proportion of ferrite was high, the tensile strength fell below 1500 MPa, and hydrogen embrittlement resistance also deteriorated. In C-3, the maximum heating temperature during the annealing process was too high, which eliminated the Mn-enriched areas. As a result, the area satisfying 1.1 × [Mnave] decreased, and hydrogen embrittlement resistance decreased. In D-3, the holding time at the maximum heating temperature in the annealing process was short, so the martensite structure ratio was less than 90% and the tensile strength was less than 1500 MPa. In E-3, the holding time at the maximum heating temperature during the annealing process was long, which caused Mn to diffuse, eliminating the dispersion of Mn-enriched areas, reducing the area satisfying 1.1 × [Mnave], and reducing hydrogen embrittlement resistance. In F-3, the cooling rate from the maximum heating temperature during the annealing process was too slow, which caused ferrite and bainite transformation, resulting in a tensile strength below 1500 MPa. The tempering temperature of I-3 was high, which reduced the dislocation density in the martensite structure, resulting in a tensile strength below 1500 MPa. In K-3, the cooling stop temperature was too high, causing bainite transformation, resulting in a tensile strength below 1500 MPa.

[0098] Fig. 1 is a graph showing the influence of the average interval of the 1.1 × [Mnave] regions and the number density of the 1.1 × [Mnave] regions on hydrogen embrittlement resistance for the steel sheets of Example 1 and Example 2. The circle in the figure indicates a steel sheet with excellent hydrogen embrittlement resistance, and the cross in the figure indicates an example with poor hydrogen embrittlement resistance. As is clear from Fig. 1, when the average interval of the 1.1 × [Mnave] regions is 10.0 µm or less and the number density of the 1.1 × [Mnave] regions is 5.0 × 10 -4 pieces / μm 2 It is clear that by controlling the above, a steel sheet with excellent hydrogen embrittlement resistance can be obtained. [Industrial Applicability]

[0099] According to the present invention, a steel sheet having high strength and excellent hydrogen embrittlement resistance can be provided. When this steel sheet is applied to automobiles, it contributes to reducing the weight of the vehicle body and improving fuel efficiency.

Claims

1. In mass%, C: 0.150-0.400%, Si: 0.01-2.00%, Mn: 0.8 to 2.0%, P: 0.0001-0.0200%, S: 0.0001-0.0200%, Al: 0.001-1.000%, N: 0.0001 to 0.0200%, O: 0.0001-0.0200%, Co: 0 to 0.500%, Ni: 0-1.000%, Mo: 0-1.000%, Cr: 0-2.000%, Ti: 0 to 0.500%, B: 0 to 0.0100%, Nb: 0 to 0.500%, V: 0 to 0.500%, Cu: 0-0.500%, W: 0-0.100%, Ta: 0-0.100%, Mg: 0 to 0.050%, Ca: 0-0.050%, Y: 0 to 0.050%, Zr: 0 to 0.050%, La: 0 to 0.050%, Ce: 0 to 0.050%, Sn: 0 to 0.050%, Sb: 0 to 0.050%, As: 0 to 0.050%, and Remainder: Fe and impurities and a chemical composition consisting of The microstructure is the area ratio, Ferrite: 5.0% or less, Martensite and tempered martensite: more than 90.0% in total, and Remainder: one or more of bainite, pearlite, and retained austenite It consists of In a cross section parallel to the sheet thickness direction and the rolling direction, when the average Mn content of the entire sheet thickness direction is [Mnave], the region where the Mn content is 1.1 × [Mnave] or more has a number density of 5.0 × 10 －４ pieces / μm ２ or more, and the distance to the nearest other region having a Mn content of 1.1 × [Mnave] or more is 10.0 μm or less on average, The tensile strength is 1500 MPa or more. A steel plate characterized by:

2. The chemical composition is Co: 0.01 to 0.500%, Ni: 0.01-1.000%, Mo: 0.01-1.000%, Cr: 0.001-2.000%, Ti: 0.001 to 0.500%, B: 0.0001 to 0.0100%, Nb: 0.001-0.500%, V: 0.001-0.500%, Cu: 0.001 to 0.500%, W: 0.001-0.100%, Ta: 0.001 to 0.100%, Mg: 0.001-0.050%, Ca: 0.001-0.050%, Y: 0.001-0.050%, Zr: 0.001 to 0.050%, La: 0.001 to 0.050%, Ce: 0.001 to 0.050%, Sn: 0.001 to 0.050%, Sb: 0.001 to 0.050%, and As: 0.001 to 0.050%, Contains one or more selected from the group consisting of The steel sheet according to claim 1 , characterized in that

3. A coating layer containing zinc, aluminum, magnesium or an alloy thereof is provided on the surface. The steel sheet according to claim 1 or 2, characterized in that

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