Steel plate

Abstract

This steel plate has a prescribed chemical composition, the microstructure thereof comprising, by area ratio, 5.0% or less ferrite and a total of greater than 90.0% martensite and tempered martensite, with the remainder being one or more of bainite, perlite, and residual austenite, and when a boundary for which the orientation difference between adjacent martensite and tempered martensite is 15 degrees or greater is defined as the prior-austenite grain boundary, the bonding strength energy EGB, as determined by the concentration of each alloy element on the prior-austenite grain boundary, is 0.50 or greater, and the tensile strength is 1500 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. 2021-172424 filed in Japan on October 21, 2021, and incorporates the content herein by reference.

Background Art

[0002] Today, with the highly specialized industrial technology fields, materials used in each technical field are required to have special and advanced performance. In particular, regarding steel sheets for automobiles, due to considerations for the global environment, in order to reduce the weight of the vehicle body and improve fuel efficiency, the demand for high-strength steel sheets has been significantly increasing. However, many metal materials deteriorate in various properties with an increase in strength, and in particular, the susceptibility to hydrogen embrittlement increases. In steel members, it is known that when the tensile strength exceeds 1200 MPa, the susceptibility to hydrogen embrittlement particularly increases, and there are cases of hydrogen embrittlement cracks in bolt steels that have been advanced in high-strength properties ahead of the automotive field. Therefore, in high-strength steel sheets with a tensile strength of 1500 MPa or more, a radical solution to hydrogen embrittlement is strongly demanded.

[0003] High-strength steel sheets with a tensile strength of 1500 MPa or more often have a microstructure mainly composed of martensite or tempered martensite. In such high-strength steel sheets, hydrogen that penetrates into the steel segregates at the grain boundaries of martensite, embrittling the grain boundaries (lowering the grain boundary strength), resulting in cracking (hydrogen embrittlement). Since hydrogen penetration occurs even at room temperature, there is no method to perfectly suppress hydrogen penetration, and modification of the internal structure of the steel is essential for a radical solution. So far, many proposals have been made for technologies to improve the hydrogen embrittlement resistance (sometimes referred to as hydrogen embrittlement resistance characteristics) of high-strength steel sheets. (For example, refer to Patent Documents 1 to 6)

[0004] In Patent Document 1, as an ultra-high strength thin steel sheet excellent in hydrogen embrittlement resistance and workability, 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 being iron and inevitable impurities, and the metallographic structure after tensile processing with a processing rate of 3% satisfying, in terms of the area ratio to the entire structure, retained austenite structure: 1% or more, bainitic ferrite and martensite: 80% or more in total, ferrite and pearlite: 9% or less (including 0%) in total, and further satisfying the average axis ratio (major axis / minor axis) of the retained austenite crystal grains: 5 or more, and a tensile strength of 1180 MPa or more, an ultra-high strength thin steel sheet excellent in hydrogen embrittlement resistance is disclosed.

[0005] In Patent Document 2, as a high strength steel sheet with a tensile strength of 1500 MPa or more, the steel composition contains Si + Mn: 1.0% or more, the main phase structure is such that ferrite and carbide form layers, and further, the aspect ratio of the carbide is 10 or more, and the interval between the said layers is 50 nm or less, and the volume ratio of the layered structure to the entire structure is 65% or more. Further, by setting the fraction of carbides having an aspect ratio of 10 or more and an angle within 25° with respect to the rolling direction among the ferrite and the carbides forming layers to 75% or more in terms of area ratio, a high strength steel sheet excellent in bendability and stress corrosion cracking resistance in the rolling direction is disclosed.

[0006] In Patent Document 3, as an ultra-high strength cold-rolled steel sheet for thin materials with excellent bendability and stress corrosion cracking resistance, it 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 inevitable impurities, and has a soft surface layer of the steel sheet satisfying the relationship of "hardness of the soft surface layer of the steel sheet / hardness of the center part of the steel sheet ≤ 0.8". The ratio of the soft surface layer of the steel sheet in the plate thickness is 0.10 or more and 0.30 or less, and the soft surface layer of the steel sheet has a tempered martensite with a volume ratio of 90% or more. The structure of the center part of the steel sheet is tempered martensite, and an ultra-high strength cold-rolled steel sheet excellent in bendability is disclosed, characterized in that the tensile strength is 1270 MPa or more.

[0007] In Patent Document 4, as a cold-rolled steel sheet with a tensile strength of 1470 MPa or more and excellent bendability and stress corrosion cracking resistance, it contains, 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 inevitable impurities, and has a metal structure with a tempered martensite phase having a volume ratio of 97% or more and a retained austenite phase having a volume ratio of less than 3%. A cold-rolled steel sheet is disclosed.

[0008] In Patent Document 5, as a super high strength steel plate with a tensile strength of 1470 MPa or more, which can exhibit excellent delayed fracture resistance even at the cut end, in terms of mass%, C: 0.15 to 0.4%, Mn: 0.5 to 3.0%, Al: 0.001 to 0.10% are each contained, the balance consists of iron and unavoidable impurities, among the unavoidable impurities, P, S, N are each limited to P: 0.1% or less, S: 0.01% or less, N: 0.01% or less, having a component composition, in terms of the area ratio to the entire structure, martensite: 90% or more, retained austenite: 0.5% or more, and a region where the local Mn concentration is 1.1 times or more the Mn content of the entire steel plate exists at 2% or more in terms of area ratio, a super high strength steel plate with a tensile strength of 1470 MPa or more is disclosed.

[0009] In Patent Document 6, as a super high strength cold rolled steel plate having excellent hydrogen embrittlement resistance and a tensile strength of 1300 MPa or more, 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%, B: 0.0001% to 0.010% are included, the values of the amount of solid solution B solB [mass%] and the prior austenite grain size Dγ [μm] satisfy the relationship of solB·Dγ≧0.0010, polygonal ferrite is 10% or less, bainite is 30% or less, retained austenite is 6% or less, tempered martensite is 60% or more, and the number density of Fe carbides in the tempered martensite is 1×10 6 / mm 2 or more, and the average dislocation density of the entire steel is 1.0×10 15 ~2.0×10 16 / m 2 , and a super high strength cold rolled steel plate having a steel structure with a crystal grain size of 7.0 μm or less is disclosed.

Prior Art Documents

Patent Documents

[0010]

Patent Document 1

[0011] As described above, several techniques for improving the hydrogen embrittlement resistance (hydrogen embrittlement resistance) of high-strength steel sheets have been proposed. However, in Patent Document 1, only the hydrogen embrittlement resistance when a stress of 1000 MPa is applied is disclosed, and no technical solution guidelines are shown for the hydrogen embrittlement resistance when a higher stress is applied. Also, as described above, hydrogen embrittlement occurs by hydrogen accumulating at grain boundaries and reducing the grain boundary bonding strength. Therefore, it is considered that if the grain boundary bonding strength can be increased, cracking due to hydrogen embrittlement can be suppressed. However, in Patent Documents 1 to 6, methods for improving the hydrogen embrittlement resistance from such a perspective have not been studied. In recent years, the requirements for hydrogen embrittlement resistance have become even more stringent, and Patent Documents 1 to 6 may not be able to meet such strict requirements. That is, conventionally, in high-strength steel sheets having a microstructure mainly composed of martensite and tempered martensite, there has been room for improvement in hydrogen embrittlement resistance.

[0012] Furthermore, in Patent Document 2, a steel sheet has a structure with pearlite as the main phase, a ferrite phase in the remaining structure having a volume ratio of 20% or less with respect to the entire structure, a lamellar spacing of the pearlite structure of 500 nm or less, and a Vickers hardness of HV200 or more. It is obtained by performing cold rolling at a rolling ratio of 60% or more (preferably 75% or more). Therefore, it can be easily estimated that the anisotropy is strong and the formability of the member by cold pressing is low. Also, in Patent Document 3, in order to improve the stress corrosion cracking characteristics, it is necessary to hold at 650 °C or 700 °C for 20 minutes or more in an atmosphere with a dew point of 15 °C or higher, and there is also a problem of low productivity.

[0013] The present invention has been made in view of the above problems. The present invention aims to provide a steel sheet having excellent hydrogen embrittlement resistance characteristics, on the premise of a high-strength steel sheet having a microstructure mainly composed of martensite and tempered martensite.

Means for Solving the Problems

[0014] As described above, hydrogen embrittlement is considered to be cracking that occurs with the grain boundary as the starting point due to the decrease in the grain boundary bonding strength caused by the segregation of hydrogen in the steel to the grain boundary. Therefore, the present inventors focused on the grain boundary bonding strength and conducted various studies on methods for improving the hydrogen embrittlement resistance characteristics. As a result, the present inventors found that by segregating a predetermined alloy element to the grain boundary, the grain boundary bonding strength is improved, and the intruded hydrogen is less likely to segregate to the grain boundary. Even if hydrogen intrudes, it is possible to suppress the decrease in the grain boundary bonding strength caused by hydrogen.

[0015] The present invention has been made in view of the above findings. The gist of the present invention is as follows. [1] The steel sheet according to one aspect of the present invention has, in mass %, C: 0.150 to 0.400%, Si: 0.01 to 2.00%, Mn: 0.80 to 2.00%, 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%, Cu: 0 to 0.500%, 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 has a microstructure which, in area ratio, is 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. When the interface with an azimuth difference of 15 deg. or more between adjacent martensite and tempered martensite is defined as the prior austenite grain boundary, the bonding strength energy E GB satisfies the following formula (1) and has a tensile strength of 1500 MPa or more. E GB = 1+(3×[Co]+0.7×[Ni]+5.5×[Mo]+0.7×[Cr]+2.9×[Ti]+47×[B]+4.3×[Nb]+4.5×[V]+5.2×[W]+3.1×[Ta]+4.3×[Zr]-0.25×[Mn]-0.1×[P]-[Cu]-1.1×[Sn]-0.6×[Sb]-0.9×[As]) ≧0.50 (1) Here, [chemical symbol] in the formula represents the concentration of each alloy element in mass % on the prior austenite grain boundary. The steel sheet described in [2][1] may contain one or more selected from the group consisting of 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: 0.001 to 0.050%, Ca: 0.001 to 0.050%, Y: 0.001 to 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%. The steel sheet described in [3][1] or [2] may have a coating layer containing zinc, aluminum, magnesium, or an alloy thereof on the surface.

Advantages of the Invention

[0016] According to the above aspect of the present invention, a steel sheet excellent in hydrogen embrittlement resistance can be provided.

Brief Description of the Drawings

[0017]

Figure 1

Embodiments for Carrying Out the Invention

[0018] Hereinafter, a steel sheet according to an embodiment of the present invention (the steel sheet according to this embodiment) will be described. The steel sheet according to this embodiment has a predetermined chemical composition, The microstructure is, in terms of area ratio, ferrite: 5.0% or less, martensite and tempered martensite: a total of more than 90.0%, and the balance: consisting of one or more of bainite, pearlite, and retained austenite, When the interface with an orientation difference of 15 deg. or more between adjacent martensite and tempered martensite is defined as the prior austenite grain boundary (prior γ grain boundary), the bonding strength energy E determined by the concentration of each alloy element on the prior austenite grain boundary GB is 0.50 or more, and the tensile strength is 1500 MPa or more.

[0019] <Chemical composition> First, the range of the content of each element constituting the chemical composition of the steel sheet according to the present embodiment will be described. Hereinafter, "%" related to the content of an element means "% by mass". In addition, the range indicated with "~" includes both end values as the lower limit or the upper limit.

[0020] C: 0.150 to 0.400% C is an element effective for increasing the tensile strength at low cost. When the C content is less than 0.150%, the target tensile strength cannot be obtained, and the fatigue characteristics of the welded part deteriorate. Therefore, the C content is set to 0.150% or more. The C content may be 0.160% or more, 0.180% or more, or 0.200% or more. On the other hand, when the C content exceeds 0.400%, the hydrogen embrittlement resistance and weldability decrease. 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.

[0021] Si: 0.01 to 2.00% Si acts as a deoxidizer and affects the form of carbides and retained austenite after heat treatment. When the Si content is less than 0.01%, it becomes difficult to suppress the formation of coarse oxides. These coarse oxides become the starting points of cracks, and the propagation of these cracks in the steel material deteriorates the hydrogen embrittlement resistance. Therefore, 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 one hand, when the Si content exceeds 2.00%, the precipitation of alloy carbides in the hot-rolled structure is delayed. 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.

[0022] Mn: 0.80 - 2.00% Mn is an element effective for increasing the strength of the steel sheet. When the Mn content is less than 0.80%, the effect cannot be obtained sufficiently. Therefore, the Mn content is set to 0.80% or more. The Mn content may be 1.00% or more, or 1.20% or more. On the other hand, when the Mn content exceeds 2.00%, Mn not only promotes the co-segregation with P and S, but may also deteriorate the corrosion resistance and hydrogen embrittlement resistance characteristics. Therefore, the Mn content is set to 2.00% or less. The Mn content may be 1.90% or less, 1.85% or less, or 1.80% or less.

[0023] P: 0.0001 - 0.0200% P is an element that strongly segregates at the ferrite grain boundaries and promotes grain boundary embrittlement. When the P content exceeds 0.0200%, the hydrogen embrittlement resistance characteristics are significantly reduced due to grain boundary embrittlement. Therefore, the P content is set to 0.0200% or less. The P content may be 0.0180% or less, 0.0150% or less, or 0.0120% or less. The lower the P content, the more preferable. However, when the P content is less than 0.0001%, the time required for refining increases, leading to a significant increase in cost. Therefore, 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.

[0024] S: 0.0001 - 0.0200% S is an element that forms non-metallic inclusions such as MnS in steel. When the S content exceeds 0.0200%, the formation of non-metallic inclusions that serve as crack initiation points during cold working becomes prominent. In this case, even if the grain boundaries are strengthened, cracks will occur from the non-metallic inclusions, and the propagation of these cracks within the steel material will deteriorate the hydrogen embrittlement resistance properties. Therefore, the S content is set to 0.0200% or less. The S content may be 0.0180% or less, 0.0150% or less, or 0.0120% or less. The lower the S content, the more preferable. However, when the S content is less than 0.0001%, the time required for refining increases, leading to a significant increase in cost. Therefore, 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.

[0025] Al: 0.001 - 1.000% Al acts as a deoxidizer for steel and is an element that stabilizes ferrite. When the Al content is less than 0.001%, the effect cannot be obtained sufficiently. 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, when the Al content exceeds 1.000%, coarse Al oxides are formed. These coarse oxides serve as crack initiation points. Therefore, when coarse Al oxides are formed, even if the grain boundaries are strengthened, cracks will occur due to the coarse oxides, and the propagation of these cracks within the steel material will deteriorate the hydrogen embrittlement resistance properties. Therefore, the Al content is set to 1.000% or less. The Al content may be 0.950% or less, 0.900% or less, or 0.800% or less.

[0026] N: 0.0001 - 0.0200% N is an element that forms coarse nitrides in the steel plate and deteriorates the hydrogen embrittlement resistance properties of the steel plate. Also, N is an element that causes the generation of blowholes during welding. If the N content exceeds 0.0200%, the hydrogen embrittlement resistance deteriorates and the generation of blow holes becomes remarkable. 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, when 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.

[0027] O: 0.0001 - 0.0200% O is an element that forms oxides and deteriorates the hydrogen embrittlement resistance. In particular, oxides often exist as inclusions. When they exist on the punched end face or the cut surface, notch-like scratches or coarse dimples are formed on the end face, which causes stress concentration during severe processing and becomes the starting point of crack formation, resulting in a significant deterioration of workability. When the O content exceeds 0.0200%, the tendency of the above-mentioned workability deterioration becomes remarkable. 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. It is preferable that the O content is less. However, setting the O content to less than 0.0001% incurs excessive cost increase and is not economically preferable. Therefore, 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.

[0028] 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 this embodiment includes the above and the balance may consist of Fe and impurities. On the other hand, for the purpose of improving various properties, the chemical composition of the steel sheet according to this embodiment may contain Co, Ni, Mo, Cr, Ti, B, Nb, V, Cu, W, Ta, Mg, Ca, Y, Zr, La, Ce, Sn, Sb, As as optional components instead of a part of the balance of Fe. Since these elements do not necessarily have to be included, the lower limit is 0%. Even if the following elements are contained as impurities, the effects of the steel sheet according to this embodiment are not inhibited.

[0029] Co: 0 to 0.500% Co is an element effective for controlling the form of carbides and increasing the strength of the steel sheet. Also, Co is an element that contributes to improving the bonding strength at grain boundaries. Therefore, Co may be contained. When sufficient effects are obtained, it is preferable that the Co content is 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, when the Co content exceeds 0.500%, coarse Co carbides precipitate. In this case, the hydrogen embrittlement resistance characteristics may deteriorate. For this reason, 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.

[0030] Ni: 0 to 1.000% Ni is an element effective for increasing the strength of the steel sheet. Also, Ni is an element that contributes to improving the bonding strength at grain boundaries. Also, Ni is an element effective for improving wettability and promoting alloying reactions. Therefore, Ni may be contained. When the above effects are obtained, it is preferable that the Ni content is 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, when the Ni content exceeds 1.000%, the hydrogen embrittlement resistance characteristics may decrease. For this reason, 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.

[0031] Mo: 0 to 1.000% Mo is an element effective in increasing the strength of the steel sheet. Also, Mo is an element having the effect of suppressing ferrite transformation that occurs during heat treatment in a continuous annealing facility or a continuous hot-dip galvanizing facility. Further, Mo is an element that also contributes to improving the bonding strength at grain boundaries. Therefore, Mo may be contained. When obtaining the above effects, it is preferable that the Mo content be 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, when the Mo content exceeds 1.000%, the effect of suppressing ferrite transformation saturates. For this reason, 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.

[0032] Cr: 0 to 2.000% Cr, like Mn, suppresses pearlite transformation and is an element effective in increasing the strength of steel. Also, Cr is an element that also contributes to improving the bonding strength at grain boundaries. Therefore, Cr may be contained. When obtaining the above effects, it is preferable that the Cr content be 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, when the Cr content exceeds 2.000%, coarse Cr carbides are formed in the center segregation part, and the hydrogen embrittlement resistance characteristics may deteriorate. For this reason, the Cr content is set to 2.000% or less. The Cr content may be 1.800% or less, 1.500% or less, or 1.000% or less.

[0033] Ti: 0 to 0.500% Ti contributes to an increase in the strength of the steel sheet through precipitation strengthening, grain refinement strengthening by suppressing the growth of ferrite crystal grains, and dislocation strengthening through suppression of recrystallization. Also, Ti is an element that also contributes to improving the bonding strength at grain boundaries. Therefore, Ti may be contained. When obtaining the above effects, it is preferable that the Ti content be 0.001% or more. The Ti content may be 0.003% or more, 0.010% or more, or 0.050% or more. On the one hand, if the Ti content exceeds 0.500%, the precipitation of carbonitrides increases, and the hydrogen embrittlement resistance characteristics 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.

[0034] B: 0 to 0.0100% B is an element that suppresses the formation of ferrite and pearlite and promotes the formation of low-temperature transformation structures such as bainite or martensite during the cooling process from the austenite temperature range. Also, B is an element beneficial for strengthening the steel. Further, B also contributes to improving the bonding strength at grain boundaries. Therefore, B may be contained. When obtaining the above effects, it is preferable that the B content is 0.0001% or more. The B content may be 0.0003% or more, 0.0005% or more, or 0.0010% or more. On the one hand, if the B content exceeds 0.0100%, coarse B oxides are formed in the steel. Since this oxide becomes the generation origin of voids during cold working, the hydrogen embrittlement resistance characteristics may deteriorate due to the formation of coarse B oxides. Therefore, the B content is set to 0.0100% or less. The B content may be 0.0080% or less, 0.0060% or less, or 0.0050% or less.

[0035] Nb: 0 to 0.500% Nb is an element effective for controlling the morphology of carbides, similar to Ti, and is also effective for improving toughness by refining the structure. Further, Nb also contributes to improving the bonding strength at grain boundaries. Therefore, Nb may be contained. When obtaining the above effects, it is preferable that the Nb content is 0.001% or more. The Nb content may be 0.002% or more, 0.010% or more, or 0.020% or more. On the one hand, if the Nb content exceeds 0.500%, the formation of coarse Nb carbides becomes significant. Since cracks are likely to occur in these coarse Nb carbides, the hydrogen embrittlement resistance characteristics may deteriorate due to the formation of coarse Nb carbides. Therefore, the Nb content is set to 0.500% or less. The Nb content may be 0.450% or less, 0.400% or less, or 0.300% or less.

[0036] V: 0 to 0.500% V contributes to the increase in the strength of the steel sheet through precipitation strengthening, grain refinement strengthening by suppressing the growth of ferrite crystal grains, and dislocation strengthening through suppression of recrystallization. Also, V contributes to the improvement of the bonding strength at grain boundaries. Therefore, V may be contained. When obtaining the above effects, it is preferable that the V content is 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, when the V content exceeds 0.500%, the precipitation of carbonitrides increases and the hydrogen embrittlement resistance characteristics may deteriorate. For this reason, 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.

[0037] Cu: 0 to 0.500% Cu is an element effective for improving the strength of the steel sheet. When the Cu content is less than 0.001%, sufficient effects cannot be obtained. For this reason, when obtaining the above effects, it is preferable that the Cu content is 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, when the Cu content exceeds 0.500%, the hydrogen embrittlement resistance characteristics may deteriorate. Also, when the Cu content is high, the steel material may embrittle during hot rolling, and in some cases, hot rolling becomes impossible. For this reason, the Cu content is set to 0.500% or less. The Cu content may be 0.450% or less, 0.400% or less, or 0.300% or less.

[0038] W: 0 to 0.100% W is an element effective for increasing the strength of the steel sheet. Also, W forms precipitates and crystallized substances. Since the precipitates and crystallized substances containing W become hydrogen trap sites, W is an element effective for improving the hydrogen embrittlement resistance characteristics. Also, W contributes to the improvement of the bonding strength at grain boundaries. Therefore, W may be contained. When obtaining the above effects, it is preferable that the W content is 0.001% or more. The W content may be 0.002% or more, 0.005% or more, or 0.010% or more. On the one hand, when the W content exceeds 0.100%, the formation of coarse W precipitates or crystallized substances becomes significant. In these coarse W precipitates or crystallized substances, cracks are likely to occur, and these cracks propagate within the steel material under a low load stress. Therefore, when coarse W precipitates or crystallized substances are formed, the hydrogen embrittlement resistance characteristics may deteriorate. For this reason, the W content is set to 0.100% or less. The W content may be 0.080% or less, 0.060% or less, or 0.050% or less.

[0039] Ta: 0 to 0.100% Ta, like Nb, V, and W, is an element effective for controlling the form of carbides and increasing the strength of the steel sheet. Also, Ta is an element that contributes to improving the bonding strength at grain boundaries. Therefore, Ta may be contained. When obtaining the above effects, it is preferable that the Ta content is 0.001% or more. The Ta content may be 0.002% or more, 0.005% or more, or 0.010% or more. On the one hand, when the Ta content exceeds 0.100%, a large number of fine Ta carbides precipitate, and with the increase in the strength of the steel sheet, the ductility may decrease, or the bending resistance and hydrogen embrittlement resistance characteristics may decrease. For this reason, the Ta content is set to 0.100% or less. The Ta content may be 0.080% or less, 0.060% or less, or 0.050% or less.

[0040] Mg: 0 to 0.050% Mg is an element that can control the form of sulfides with a trace content. Therefore, Mg may be contained. When obtaining the above effects, it is preferable that the Mg content is 0.001% or more. The Mg content may be 0.005% or more, 0.010% or more, or 0.020% or more. On the one hand, when the Mg content exceeds 0.050%, coarse inclusions may be formed and the hydrogen embrittlement resistance characteristics may decrease. For this reason, the Mg content is set to 0.050% or less. The Mg content may be 0.040% or less, 0.030% or less, or 0.020% or less.

[0041] Ca: 0 to 0.050% Ca is not only useful as a deoxidizing element but also an element effective in controlling the morphology of sulfides. Therefore, Ca may be contained. When obtaining 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, when the Ca content exceeds 0.050%, coarse inclusions may be formed and the hydrogen embrittlement resistance characteristics may decrease. For this reason, 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.

[0042] Y: 0 to 0.050% Y is an element that can control the morphology of sulfides with a trace amount of content, similar to Mg and Ca. Therefore, Y may be contained. When obtaining 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, when the Y content exceeds 0.050%, coarse Y oxides may be generated and the hydrogen embrittlement resistance characteristics may decrease. For this reason, 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.

[0043] Zr: 0 to 0.050% Zr is an element that can control the morphology of sulfides with a trace amount of content, similar to Mg, Ca, and Y. Also, Zr is an element that contributes to improving the bonding strength at grain boundaries. Therefore, Zr may be contained. When obtaining 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, when the Zr content exceeds 0.050%, coarse Zr oxides may be generated and the hydrogen embrittlement resistance characteristics may decrease. For this reason, the Zr content is set to 0.050% or less. The Zr content may be 0.040% or less, 0.030% or less, or 0.020% or less.

[0044] La: 0 to 0.050% La, like Mg, Ca, Y, and Zr, is an element that can control the form of sulfides even with a trace amount of content. Therefore, La may be contained. When obtaining 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, when the La content exceeds 0.050%, La oxide may be generated, and the hydrogen embrittlement resistance characteristics may decrease. 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.

[0045] Ce: 0 to 0.050% Ce, like La, is an element that can control the form of sulfides even with a trace amount of content. Therefore, Ce may be contained. When obtaining 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, when the Ce content exceeds 0.050%, Ce oxide may be generated, and the hydrogen embrittlement resistance characteristics may decrease. Therefore, the Ce content is set to 0.050% or less. The Ce content may be 0.040% or less, 0.030% or less, or 0.020% or less.

[0046] 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 characteristics may decrease due to grain boundary embrittlement. This adverse effect becomes particularly prominent when the Sn content exceeds 0.050%. Therefore, the Sn content is set to 0.050% or less. The Sn content may be 0.040% or less, 0.030% or less, or 0.020% or less. The lower the Sn content, the more preferable it is, and it may be 0%, but when 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.

[0047] Sb: 0 to 0.050% Sb is an element contained when using scrap as a steel raw material, similar to Sn. Sb is an element that strongly segregates at grain boundaries, leading to embrittlement of grain boundaries and a decrease in ductility. This adverse effect becomes particularly prominent when the Sb content exceeds 0.050%. Therefore, the Sb content is set to 0.050% or less. The Sb content may be 0.040% or less, 0.030% or less, or 0.020% or less. The lower the Sb content, the more preferable it is, and it may even be 0%. However, when 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.

[0048] As: 0 - 0.050% As is an element contained when using scrap as a steel raw material, similar to Sn and Sb. It strongly segregates at grain boundaries, leading to embrittlement of grain boundaries and a decrease in ductility. When the As content is high, the hydrogen embrittlement resistance characteristics may decrease. This adverse effect becomes particularly prominent when the As content exceeds 0.050%. Therefore, the As content is set to 0.050% or less. The As content may be 0.040% or less, 0.030% or less, or 0.020% or less. The lower the As content, the more preferable it is, and it may even be 0%. However, when 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.

[0049] As described above, the chemical composition of the steel sheet according to this embodiment may contain basic components, and the balance may consist of Fe and impurities. It may contain basic components, further contain one or more optional components, and the balance may consist of Fe and impurities.

[0050] The chemical composition of the steel sheet according to this embodiment may be measured by a general method. For example, it may be measured using ICP-AES (Inductively Coupled Plasma-Atomic Emission Spectrometry) for chips in accordance with JIS G1201:2014. In this case, the chemical composition is the average content over the entire plate thickness. For C and S that cannot be measured by ICP-AES, combustion-infrared absorption method may be used, for N, inert gas fusion-thermal conductivity method may be used, and for O, inert gas fusion-non-dispersive infrared absorption method may be used for measurement. When the steel sheet is provided with a coating layer on its surface, the coating layer may be removed by mechanical grinding or the like and then the chemical composition may be analyzed. When the coating layer is a plating layer, it may also be removed by dissolving the plating layer in an acid solution added with an inhibitor for suppressing corrosion of the steel sheet.

[0051] <Microstructure (metallic structure)> Next, the microstructure of the steel sheet according to this embodiment will be described. In this embodiment, the microstructure is the microstructure at a position in the range of 1 / 8 to 3 / 8 of the plate thickness (t / 4 part) in the plate thickness direction from the surface of the steel sheet. The reason for defining the microstructure of the t / 4 part is that it is a typical microstructure of the steel sheet and has a large correlation with the properties of the steel sheet. Also, unless otherwise specified, the fraction (%) of each of the following phases is the area ratio.

[0052] Ferrite: 5.0% or less Ferrite affects the formability of steel with martensite as the main structure. As the area ratio of ferrite increases, the local formability and hydrogen embrittlement resistance properties decrease. In particular, when the area ratio of ferrite exceeds 5.0%, the hydrogen embrittlement resistance properties may decrease due to fracture in elastic deformation during stress loading. Therefore, the area ratio of ferrite is set to 5.0% or less. The area ratio of ferrite may 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% requires advanced control in manufacturing and causes a decrease in yield. Therefore, the area ratio of ferrite may be 1.0% or more.

[0053] Martensite and tempered martensite: more than 90.0% in total The total area ratio of martensite and tempered martensite affects the strength of the steel, and the greater the area ratio, the higher the tensile strength. If the total area ratio of martensite and tempered martensite is 90.0% or less, the target tensile strength cannot be achieved. In addition, it may cause fracture during elastic deformation under stress loading, or may cause a decrease in hydrogen embrittlement resistance due to the formation of structures other than martensite and tempered martensite and an increase in the inhomogeneity of the microstructure. Therefore, 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%.

[0054] Remainder: composed 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 when the 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%, it may cause fracture during elastic deformation under stress loading and reduce the hydrogen embrittlement resistance. Therefore, the area ratio of the remaining structure is preferably 8.0% or less, more preferably 7.0% or less. Among these, in particular, pearlite and retained austenite are structures that deteriorate the local ductility of the steel, and the less the better. On the other hand, setting the area ratio of the remaining structure to 0% requires advanced control in manufacturing, which may lead to a decrease in yield. Therefore, the area ratio of the remaining structure may be 1.0% or more.

[0055] The area ratio of each phase in the microstructure of the steel sheet according to this embodiment can be obtained by the following method.

[0056] (Evaluation method for the area ratio of ferrite) The area ratio of ferrite is determined by observing the t / 4 part (the range of 1 / 8 to 3 / 8 of the plate thickness centered at the position of 1 / 4 of the plate thickness in the plate thickness direction from the surface) using an electron channeling contrast image obtained by a field emission-scanning electron microscope (FE-SEM). The electron channeling contrast image is a technique for detecting the crystal orientation difference within a crystal grain as a difference in image contrast. In this image, the portion that appears with a uniform contrast in the structure determined to be ferrite rather than pearlite, bainite, martensite, or retained austenite is polygonal ferrite. The area ratio of polygonal ferrite in each of 8 fields of an electron channeling contrast image of 35 μm × 25 μm is calculated by an image analysis method, and the average value thereof is taken as the area ratio of ferrite.

[0057] (Method for evaluating the total area ratio of martensite and tempered martensite) The total area ratio of martensite and tempered martensite is also determined from the image taken by the aforementioned electron channeling contrast. Since these structures are less likely to be etched than ferrite, they exist as convex portions on the tissue observation surface. Tempered martensite is an aggregate of lath-shaped crystal grains and contains iron-based carbides with a major axis of 20 nm or more inside, and the carbides belong to a plurality of variants, that is, a plurality of groups of iron-based carbides extending in different directions. Also, retained austenite exists as a convex portion on the tissue observation surface. Therefore, by subtracting the area ratio of the convex portion obtained by the above procedure from the area ratio of retained austenite measured by the procedure described later, it becomes possible to accurately measure the total area ratio of martensite and tempered martensite.

[0058] (Method for evaluating the total area ratio of bainite, pearlite, and retained austenite) The area ratio of retained austenite can be calculated by measurement using X-rays. That is, from the surface of the sample plate to a position of 1 / 4 of the plate thickness in the plate thickness direction is removed by mechanical polishing and chemical polishing. Then, from the integrated intensity ratio of the diffraction peaks of (200), (211) of the bcc phase and (200), (220), (311) of the fcc phase obtained using MoKα rays as characteristic X-rays for the polished sample, the tissue fraction of retained austenite is calculated, and this is taken as the area ratio of retained austenite. Also, the area ratio of pearlite is determined from the image taken by the above-described electron channeling contrast. Pearlite is a structure in which plate-like carbides and ferrite are arranged side by side. Also, bainite is an aggregate of lath-shaped crystal grains that does not contain iron-based carbides with a major axis of 20 nm or more inside, or contains iron-based carbides with a major axis of 20 nm or more inside, and the carbides belong to a single variant, that is, an iron-based carbide group extending in the same direction. Here, the iron-based carbide group extending in the same direction means that the difference in the extension direction of the iron-based carbide group is within 5°.

[0059] <Bonding strength energy> Cracks due to hydrogen embrittlement occur when hydrogen in steel segregates at grain boundaries, resulting in a decrease in the bonding strength of the grain boundaries, and the grain boundaries with decreased bonding strength serve as the starting points for cracks. On the other hand, when a predetermined alloy element is segregated at the grain boundaries, the bonding strength of the grain boundaries is improved. Also, due to the already segregated alloy elements, it becomes difficult for the intruding hydrogen to segregate at the grain boundaries, and even if hydrogen intrudes, it is possible to suppress the decrease in the bonding strength of the grain boundaries due to hydrogen. The present inventors investigated the contribution degree to the improvement of grain boundary strength by each alloy element, taking the prior austenite grain boundaries, which are interfaces (martensite / tempered martensite, martensite / martensite, or tempered martensite / tempered martensite interfaces) where the orientation difference between adjacent martensite and tempered martensite is 15 deg. or more, as the main grain boundaries in a microstructure mainly composed of martensite and tempered martensite. As a result, the bonding strength energy E GB can be expressed as in the following formula (1) using the concentration of each alloy element on the grain boundary, EGB It has been found that when it is 0.50 or more, the hydrogen embrittlement resistance property is clearly improved. Therefore, in the steel sheet according to the present embodiment, when the interface with an orientation difference of 15 deg. or more between adjacent martensite and tempered martensite is defined as the prior austenite grain boundary, the bonding strength energy E determined by the concentration of each alloy element on the prior austenite grain boundary GB satisfies the following formula (1). E GB =1+(3×[Co]+0.7×[Ni]+5.5×[Mo]+0.7×[Cr]+2.9×[Ti]+47×[B]+4.3×[Nb]+4.5×[V]+5.2×[W]+3.1×[Ta]+4.3×[Zr]-0.25×[Mn]-0.1×[P]-[Cu]-1.1×[Sn]-0.6×[Sb]-0.9×[As]) ≧0.50 (1) Here, [chemical symbol] in the formula represents the concentration of each alloy element in mass % on the prior austenite grain boundary. The reason for targeting the interface with an orientation difference of 15 deg. or more is that hydrogen has a tendency to preferentially accumulate at the prior austenite grain boundary of 15 deg. or more. As can be seen from formula (1), not all of the alloy elements segregating at the grain boundary increase the bonding strength energy. By segregating many alloy elements that increase the grain boundary bonding energy, the grain boundary bonding energy becomes high. Here, the interface with an orientation difference of 15 deg. or more between adjacent martensite and tempered martensite refers to the interface between martensite and martensite with an orientation difference of 15 deg. or more, the interface between martensite and tempered martensite with an orientation difference of 15 deg. or more, and the interface between tempered martensite and tempered martensite with an orientation difference of 15 deg. or more.

[0060] The concentration of each alloy element on the prior austenite grain boundary is determined by observing the t / 4 part (the range of 1 / 8 to 3 / 8 of the plate thickness centered at the position of 1 / 4 of the plate thickness in the plate thickness direction from the surface) in the same manner as the above SEM observation using the EDS (Energy Dispersive X-ray Spectrometer) of the TEM (Transmission Electron Microscope). More specifically, as the TEM (Transmission Electron Microscope), a Cs-corrected TEM (Spherical Aberration Corrected TEM) is used. The thin film sample used for TEM observation is obtained by the following method. A sample for measuring the alloy element amount is collected from the range of 1 / 8 to 3 / 8 of the plate thickness from the steel plate sample, and wet polishing using emery paper is performed until it is polished to a thickness of about 100 μm. After that, electrolytic polishing is performed until the thickness at which TEM observation is possible by twin jet electrolytic polishing. The electrolytic polishing method is performed using a twin jet electrolytic polishing device. Since the appropriate conditions for twin jet electrolytic polishing vary depending on the base material component of the sample, it is necessary to extract for each sample. After performing twin jet, the quantification accuracy of the elements on the prior austenite grain boundary is improved by uniformly milling the thin film sample using Ar ion milling. The thin film sample thus obtained is observed with a Cs-corrected TEM. The observation position is on the prior austenite grain boundary, and the prior austenite grain boundary is found as follows. In the bright-field image (Bright-Field image) in TEM observation, when observed at 30,000 times magnification, the prior austenite grain boundary, packet boundary, and block boundary appear as black lines. Among these black lines, the sample is tilted and rotated so that the black line indicating any prior austenite grain boundary is horizontal to the incident direction of the electron beam of the TEM. In that state, elemental analysis using EDS is performed at 100,000 times magnification directly above the prior austenite grain boundary. The elemental analysis using the integration count of EDS analysis is performed by the following method. Point analysis is performed 3 times directly above the prior austenite grain boundary to quantify the alloy element concentration on the prior austenite grain boundary. This analysis is performed on 5 prior austenite grain boundaries, and the average alloy element concentration is calculated. This average alloy element concentration is taken as the alloy element concentration on the prior austenite grain boundary.

[0061] (Mechanical properties) In the steel sheet according to this embodiment, the tensile strength (TS) is 1500 MPa or more as the strength contributing to the weight reduction of the automobile body. Although it is not necessary to limit the upper limit, when the tensile strength increases, the formability may decrease, so the tensile strength may be 2000 MPa or less.

[0062] (Thickness) The steel sheet according to this embodiment is not limited in terms of thickness, but is preferably 1.0 to 2.2 mm. More preferably, the thickness is 1.05 mm or more, and even more preferably 1.1 mm or more. Also, more preferably, the thickness is 2.1 mm or less, and even more preferably 2.0 mm or less.

[0063] (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. This coating layer may be composed of zinc, aluminum, magnesium, or an alloy thereof and impurities. By providing a coating layer on the surface, the corrosion resistance is improved. In the case of a steel sheet for automobiles, if there is a concern about pitting due to corrosion, even if the strength is increased, it may not be possible to reduce the thickness below a certain thickness. Since one of the purposes of increasing the strength of the steel sheet is weight reduction by reducing the thickness, even if a high-strength steel sheet is developed, if the corrosion resistance is low, the applicable parts are limited. As a method for solving these problems, it is conceivable to form a coating layer on the front and back surfaces to improve the corrosion resistance. Even if a coating layer is formed, the hydrogen embrittlement resistance characteristics of the steel sheet according to this embodiment are not impaired. The coating layer is, for example, a hot-dip zinc coating layer, an alloyed hot-dip zinc coating layer, an electro-galvanized coating layer, an aluminum coating layer, a Zn-Al alloy coating layer, an Al-Mg alloy coating layer, or a Zn-Al-Mg alloy coating layer. When having a coating layer on the surface, the surface serving as the reference for the t / 4 part described above is the surface of the base metal excluding the coating layer.

[0064] (Manufacturing method) The steel sheet according to this embodiment can obtain the above effects as long as it has the above characteristics regardless of the manufacturing method, but it can be manufactured by a manufacturing method including the following steps (I) to (VII). (I) Heating step of heating a steel slab having a predetermined chemical composition (II) Hot rolling step of hot rolling the heated steel slab to obtain a hot rolled steel sheet (III) Cooling step of starting cooling within 3.0 seconds after completion of the hot rolling step and cooling the hot rolled steel sheet to a coiling temperature of 550 to 700 °C at an average cooling rate of 20 °C / second or more and 50 °C / second or less (IV) Coiling step of coiling the hot rolled steel sheet after the cooling step at the coiling temperature (V) Holding step of holding the hot rolled steel sheet after the coiling step in a temperature range of 400 to 550 °C for 600 seconds or more (VI) Cold rolling step of pickling and cold rolling the hot rolled steel sheet after the holding step to obtain a cold rolled steel sheet (VII) Annealing step of annealing the cold rolled steel sheet after the cold rolling step by holding it at an annealing temperature of 800 °C or more and less than 900 °C Hereinafter, the preferable conditions in each step will be described.

[0065] (Heating step) In the heating step, a steel slab or the like having the same chemical composition as the steel sheet according to this embodiment is heated prior to hot rolling. The heating temperature is not limited as long as the rolling temperature in the next step can be ensured. For example, it is 1000 to 1300 °C. The steel slab to be used is preferably cast by the continuous casting method from the viewpoint of productivity, but it may be manufactured by the ingot method or the thin slab casting method. When the steel slab obtained by continuous casting can be directly supplied to the hot rolling step at a sufficiently high temperature, the heating step may be omitted.

[0066] (Hot rolling step) In the hot rolling step, the heated steel slab is hot rolled to obtain a hot rolled steel sheet. The hot rolling process includes rough rolling and finish rolling. In the finish rolling, rolling reduction is performed in multiple passes. Among the multiple passes, 4 or more passes are large reduction passes with a rolling reduction rate of 20% or more, and the inter-pass time for each of the large reduction passes is set to 5.0 seconds or less. Also, the rolling start temperature is set to 950 - 1100 °C, and the rolling end temperature is set to 800 - 950 °C.

[0067] [In finish rolling, large reduction passes with a rolling reduction rate of 20% or more: 4 or more passes] [Inter-pass time: within 5.0 seconds] By controlling the rolling reduction rate, number of rolling passes, and inter-pass time in finish rolling, it becomes possible to control the morphology of austenite grains to be equiaxed and fine. When the austenite grains become equiaxed and fine, the grain boundary diffusion of alloying elements is promoted, and the precipitation of alloy carbides or nitrides at grain boundaries is promoted. If the number of passes with a rolling reduction rate of 20% or more (large reduction passes) is less than 4 passes, unrecrystallized austenite remains, and thus sufficient effects cannot be obtained. Therefore, in 4 or more passes, the rolling reduction rate is set to 20% or more (rolling reduction is performed 4 or more times with a rolling reduction rate of 20% or more). Preferably, in 5 or more passes, the rolling reduction rate is set to 20% or more. On the other hand, although there is no particular limitation on the upper limit of the number of passes with a rolling reduction rate of 20% or more, if it exceeds 10 passes, it is necessary to install a large number of rolling stands, which may lead to an increase in equipment size and manufacturing cost. Therefore, the number of passes (number of passes) with a rolling reduction rate of 20% or more may be 10 passes or less, 9 passes or less, or 7 passes or less. Also, the inter-pass time in finish rolling has a great influence on the recrystallization and grain growth of austenite grains after rolling. Even when the number of large reduction passes is 4 or more, if the inter-pass time for each pass exceeds 5.0 seconds, grain growth is likely to occur, and the austenite grains become coarser. On the other hand, although it is not necessary to limit the lower limit of the inter-pass time, if the inter-pass time for each of the large reduction passes is less than 0.2 seconds, the recrystallization of austenite is not completed, the proportion of unrecrystallized austenite increases, and sufficient effects may not be obtained. Therefore, it is preferable to set the inter-pass time for the large reduction passes to 0.2 seconds or more. The inter-pass time may be 0.3 seconds or more, or 0.5 seconds or more.

[0068] (Cooling process) In the cooling process, the hot-rolled steel sheet after the hot rolling process is cooled to a coiling temperature of 550 to 700°C at an average cooling rate of 20°C / sec or more and 50°C / sec or less, starting the cooling within 3.0 seconds from the completion of the hot rolling process (the completion of the final pass of the finish rolling). If the time from the completion of hot rolling to the start of rolling exceeds 3.0 seconds, or the average cooling rate to the coiling temperature is less than 20°C / sec, ferrite transformation occurs from austenite before coiling. In this case, the driving force for precipitation becomes small, and it becomes difficult for precipitates to precipitate uniformly and finely in the subsequent process. On the other hand, if the average cooling rate to the coiling temperature is too fast, a hardening phase is likely to be formed. In this case, the subsequent manufacturability is significantly deteriorated and the productivity is reduced. Also, if the cooling stop temperature is less than 550°C, the precipitation of precipitates is delayed, leading to deterioration of manufacturability and a decrease in productivity. Further, if the cooling stop temperature exceeds 700°C, ferrite transformation occurs from austenite, the driving force for carbide precipitation becomes small, and it becomes difficult for precipitates to precipitate uniformly and finely in the subsequent process, which is not preferable. Also, if the cooling stop temperature exceeds 700°C, an internal oxidation layer is likely to be formed on the steel sheet surface, cracks are likely to occur on the surface, or the productivity in the subsequent pickling process is significantly deteriorated, which is not preferable. There is no need to limit the lower limit of the time from the completion of hot rolling to the start of rolling, and it may be carried out in as short a time as possible within the range of equipment limitations.

[0069] (Coiling process) (Soaking process) In the coiling process, the hot-rolled steel sheet after the cooling process is coiled at the coiling temperature (cooling stop temperature). Further, in the subsequent soaking process, the coiled hot-rolled steel sheet is held (soaked) in the temperature range of 400 to 550°C for 600 seconds or more. By controlling these coiling and soaking conditions, alloy carbides or nitrides are precipitated on the steel sheet. The precipitates deposited here can be made to be unevenly distributed at the prior austenite grain boundaries by controlling the subsequent process. If the holding temperature is too high, the precipitates will coarsen and will not disperse uniformly. Also, if the holding temperature is too low, although the precipitates will become finer, it will take a long time to complete precipitation, resulting in a decrease in manufacturability and productivity. Also, if the holding time is short, the alloy carbides will not precipitate sufficiently. To hold under the above conditions, for example, methods such as covering with a cover or coating with a heating box may be employed.

[0070] (Cold rolling process) In the cold rolling process, the hot-rolled steel sheet after the holding process is unwound, pickled, and cold-rolled to obtain a cold-rolled steel sheet. By performing pickling, the oxide scale on the surface of the hot-rolled steel sheet can be removed, thereby improving the chemical conversion treatability and plating property of the cold-rolled steel sheet. Pickling may be carried out under known conditions, either once or in multiple steps. The reduction ratio (rolling ratio) of cold rolling is not particularly limited. For example, it is 20 to 80%.

[0071] (Annealing process) In the annealing process, the cold-rolled steel sheet after the cold rolling process is held and annealed at an annealing temperature of 800°C or higher and less than 900°C. In this annealing process, in the process of heating to the annealing temperature which is the austenite single-phase region, in a relatively low temperature range, these precipitates play a role of pinning the prior austenite grain boundaries. As a result, the precipitates are unevenly distributed on the prior austenite grain boundaries. Further heating is performed, and when reaching a relatively high temperature range, the precipitates become thermally unstable and dissolve. As a result, the alloy elements can be segregated on the prior austenite grain boundaries. In this state, as described in the post-annealing cooling process to be described later, by performing rapid cooling, austenite transforms into martensite, and a high-strength steel sheet with alloy elements unevenly distributed on the prior austenite grain boundaries can be obtained. If the annealing temperature is less than 800°C, the amount of austenite formed is small, and the dissolution of carbides is insufficient. Therefore, the annealing temperature is set to 800°C or higher. The annealing temperature is preferably 830°C or higher. On the one hand, when the annealing temperature is 900 °C or higher, grain growth occurs and the grain size of the prior austenite coarsens, suppressing the segregation of a predetermined alloy element to the grain boundaries and possibly deteriorating the hydrogen embrittlement resistance characteristics. The holding time at the annealing temperature does not need to be limited. However, if the holding time is less than 10 seconds, the fraction of austenite at the annealing temperature may be insufficient or the dissolution of carbides may be insufficient. Therefore, the holding time is preferably 10 seconds or more. On the other hand, even if the holding time is long, there is no problem in terms of characteristics. However, assuming continuous annealing, the line length of the equipment becomes long, so about 600 seconds may be set as the substantial upper limit. The average heating rate up to the annealing temperature is preferably 2 to 35 °C / second.

[0072] (Post-annealing cooling process) (Tempering process) In the post-annealing cooling process, after the annealing process, the cold-rolled steel sheet may be cooled from the annealing temperature to 25 °C to 300 °C at an average cooling rate of 20 to 100 °C / second. By this cooling, the steel sheet is rapidly cooled in a state where alloy elements are segregated at the austenite grain boundaries, and the austenite transforms into martensite. As a result, a steel sheet having a structure mainly composed of martensite with alloy segregation at the prior austenite grain boundaries is obtained. If the average cooling rate is less than 20 °C / second, a sufficient amount of martensite is not generated. On the other hand, when the average cooling rate exceeds 100 °C / second, assuming continuous annealing, the equipment capacity may be insufficient and equipment enhancement may be required. Therefore, 100 °C / second is set as the substantial upper limit. Also, if the cooling stop temperature exceeds 300 °C, the untransformed austenite that has not undergone martensite transformation is likely to undergo bainite transformation, and the strength may decrease. On the other hand, even if the cooling stop temperature is set to less than 25 °C, the effect saturates, and a special refrigerant or the like is required, resulting in a decrease in productivity or an increase in cost. The cold-rolled steel sheet after the annealing post-cooling process may further be subjected to a tempering process of heating to 50 to 550 °C and holding for 10 to 1000 seconds. By performing this tempering, alloying elements that could not fully segregate at the grain boundaries and are present within the grains can be segregated onto the prior austenite grain boundaries. Also, by tempering martensite into tempered martensite, formability can be improved. If the tempering temperature (holding temperature) is less than 50 °C or the holding time is less than 10 seconds, the above effects cannot be obtained. On the other hand, if the tempering temperature exceeds 550 °C, a decrease in strength may occur due to a decrease in the dislocation density in the tempered martensite, leading to a decrease in tensile strength. Also, carbides may coarsely precipitate on the prior austenite grain boundaries, sometimes deteriorating the hydrogen embrittlement resistance characteristics. Further, if the holding time exceeds 1000 seconds, not only does the strength decrease, but productivity also decreases. Tempering may be carried out in a continuous annealing facility or may be performed offline after continuous annealing using separate equipment.

[0073] In the above annealing post-cooling process, during cooling, holding may be carried out for 10 to 200 seconds in a temperature range of 350 to 650 °C (the second temperature range: a temperature range considered to be higher than the Ms point). In this case, the cooling rate up to the second temperature range excluding the holding and the average cooling rate from the holding temperature to 25 to 300 °C (cooling stop temperature) may each be 20 to 100 °C / second. That is, in this case, after the annealing process, cooling is carried out from the annealing temperature at an average cooling rate of 20 to 100 °C / second to the second temperature range of 350 to 650 °C, holding is carried out for 10 to 200 seconds in the second temperature range, and cooling is carried out from the second temperature range at an average cooling rate of 20 to 100 °C / second to 25 to 300 °C. By holding at 350 to 650 °C, alloying elements that cannot be completely segregated at grain boundaries and are present within grains can be segregated onto prior austenite grain boundaries, improving the hydrogen embrittlement resistance characteristics. However, if the holding temperature is less than 350 °C, bainite transformation is likely to occur, possibly reducing the strength. If the holding time is less than 20 seconds, the effect of segregating the elements present within grains onto prior austenite grain boundaries cannot be obtained. On the other hand, if the holding temperature exceeds 650 °C, ferrite transformation from austenite is likely to occur, leading to a decrease in tensile strength. Also, if the holding time exceeds 200 seconds, bainite transformation or ferrite transformation from austenite is likely to occur. The range of the holding temperature is preferably 370 °C or higher and 630 °C or lower, more preferably 390 °C or higher and 610 °C or lower. The range of the holding time is preferably 30 seconds or longer and 180 seconds or shorter, more preferably 50 seconds or longer and 160 seconds or shorter. Holding during the intermediate cooling process after annealing and the tempering process both further promote the segregation of alloying elements onto prior austenite grain boundaries. Therefore, either one of the processes may be performed, or both processes may be performed. It is not necessary to perform either process.

[0074] In the method for manufacturing a steel sheet according to the present embodiment, a coating layer forming step of forming a coating layer on the surface (one or both sides) of the steel sheet may be provided. As the coating layer, a coating layer containing zinc, aluminum, magnesium, or an alloy thereof is preferable. The coating layer is, for example, a plating layer. The coating method is not limited, but for example, when forming a zinc-based coating layer by hot dip plating, the cold-rolled steel sheet is adjusted (heated or cooled) so that the steel sheet temperature becomes (plating bath temperature - 40) °C to (plating bath temperature + 50) °C, and then immersed in a plating bath at 450 to 490 °C to form a plating layer. This is an example of the conditions. The reason why this condition is preferable is that if the steel sheet temperature during immersion in the plating bath is lower than the molten zinc plating bath temperature - 40 °C, the heat extraction during immersion in the plating bath is large, and a part of the molten zinc may solidify, possibly deteriorating the plating appearance. If it exceeds the molten zinc plating bath temperature + 50 °C, it will induce operational problems associated with the increase in the plating bath temperature. When forming a plating layer mainly composed of zinc, the composition of the plating bath is such that the effective Al amount (the value obtained by subtracting the total Fe amount from the total Al amount in the plating bath) is 0.050 to 0.250% by mass, and preferably contains Mg as required, with the balance being Zn and impurities. When the effective Al amount in the plating bath is less than 0.050% by mass, the intrusion of Fe into the plating layer may proceed excessively, and the plating adhesion may decrease. On the other hand, when the effective Al amount in the plating bath exceeds 0.250% by mass, an Al-based oxide that inhibits the movement of Fe atoms and Zn atoms may be generated at the boundary between the steel sheet and the plating layer, and the plating adhesion may decrease.

[0075] The formation of the above coating layer may be carried out after the annealing post-cooling process described above, or may be carried out during the annealing post-cooling process or during the tempering process. That is, it may be carried out as part of the holding at 350 to 650°C during the annealing post-cooling process, or as part of the holding at 50 to 550°C during the tempering process.

[0076] When forming a plating layer mainly composed of zinc as the coating layer, alloying treatment may be further carried out. In this case, conditions such as holding the steel sheet on which the plating layer is formed at 480 to 550°C for 1 to 30 seconds are exemplified. The alloying process may also be carried out during the annealing post-cooling process or during the tempering process described above. That is, it may be carried out as part of the holding at 350 to 650°C during the annealing post-cooling process, or as part of the holding at 50 to 550°C during the tempering process.

[0077] On the surface of the coating layer, upper-layer plating may be applied for the purpose of improving paintability and weldability, or various treatments such as chromate treatment, phosphate treatment, lubricity improvement treatment, weldability improvement treatment, etc. may be carried out.

Examples

[0078] Examples of the present invention are shown below. The examples shown below are examples of the present invention, and the present invention is not limited to the examples described below.

[0079] <Example 1> Steel having the chemical compositions shown in Tables 1-1 to 1-4 was melted and steel slabs were cast. This steel slab was inserted into a furnace heated to 1220°C, held for 60 minutes, then taken out into the atmosphere and hot-rolled to obtain a steel plate with a thickness of 2.8 mm (hot-rolled steel plate). In the hot rolling, a rolling mill having seven stands was used, and a total of seven finishing passes were continuously performed (so that the inter-pass time was constant), among which, four rolling passes with a reduction rate exceeding 20% were given. Also, the inter-pass time between each rolling pass with a reduction rate of 20% or more in the finishing passes and the rolling pass immediately before each such rolling pass was set to 0.6 seconds. The starting temperature of the finishing rolling was 1060°C and the ending temperature was 870°C. For this hot-rolled steel plate, after the hot rolling was completed, cooling was started by water cooling 2.2 seconds later, cooled to 580°C at an average cooling rate of 38.0°C / second and then coiled, and then charged into a furnace at 530°C and held for 1800 seconds. Subsequently, the oxide scale of this hot-rolled steel plate was removed by pickling, and cold rolling with a reduction rate of 50.0% was performed to obtain a cold-rolled steel plate with a thickness of 1.4 mm. This cold-rolled steel plate was heated to 880°C at an average heating rate of 12.0°C / second, held at 880°C for 120 seconds, and then cooled to 150°C at an average cooling rate of 42.0°C / second. Thereafter, the cold-rolled steel plate was reheated to 230°C and subjected to tempering by holding for 180 seconds. No plating treatment was performed.

[0080] The chemical composition was analyzed using samples taken from the obtained steel plate. As a result, it was equivalent to the chemical compositions of the steels shown in Tables 1-1 to 1-4.

[0081] Also, for the obtained cold-rolled steel plate, the area ratios of ferrite, martensite and tempered martensite, and the balance (one or more of bainite, pearlite and retained austenite) in the microstructure of the t / 4 part were determined by the method described above. The results are shown in Table 2.

[0082] Also, by the method described above, the concentration of each alloy element on the austenite grain boundary was measured, and E GB was determined.

[0083] In addition, the tensile strength, total elongation, and hydrogen embrittlement resistance (hydrogen embrittlement resistance characteristics) of the obtained cold-rolled steel sheet were evaluated by the following methods.

[0084] (Method for evaluating tensile properties) The tensile test was conducted in accordance with JIS Z 2241 (2011). A JIS No. 5 test piece was taken from the direction in which the longitudinal direction of the test piece was parallel to the direction perpendicular to the rolling direction of the steel strip, and the tensile strength (TS) and total elongation (El) were measured.

[0085] (Method for evaluating hydrogen embrittlement resistance) Regarding the hot-dip galvanized steel sheet manufactured using the method for manufacturing a steel sheet according to an embodiment of the present invention, the hydrogen embrittlement resistance was evaluated by the following method. Specifically, after shearing the steel sheet with a clearance of 15%, a U-bending test was performed at 8R. A strain gauge was attached to the center of the obtained test piece, and stress was applied by tightening both ends of the test piece with bolts. The applied stress was calculated from the strain of the monitored strain gauge. The load stress was set to a stress corresponding to 80% of the tensile strength (TS) (for example, in the case of A in Table 2, the applied stress = 1515 MPa × 0.8 = 1212 MPa). This is because it is considered that the residual stress introduced during forming corresponds to the tensile strength of the steel sheet. The obtained U-bending test piece was immersed in an HCl aqueous solution with a liquid temperature of 25°C and a pH of 2, and held for 96 hours to check for cracks. Since the lower the pH of the HCl aqueous solution and the longer the immersion time, the larger the amount of hydrogen that penetrates into the steel sheet, the hydrogen embrittlement environment becomes severe. After immersion, when a crack with a length exceeding 1.00 mm was observed in the U-bending test piece, it was evaluated as NG, and when no crack with a length exceeding 1.00 mm was observed, it was evaluated as OK. When the evaluation was OK, it was considered qualified, and when it was NG, it was considered unqualified.

[0086] When the tensile strength was 1500 MPa or more and the evaluation of hydrogen embrittlement resistance was OK, it was evaluated as a steel sheet with high strength and excellent hydrogen embrittlement resistance.

[0087]

Table 1-1

[0088]

Table 1-2

[0089]

Table 1-3

[0090]

Table 1-4

[0091]

Table 2

[0092] Referring to Table 1-1 to Table 2, since the C content of Production No. P was low, the tensile strength was less than 1500 MPa. Since the C content of Production No. Q was high, the hydrogen embrittlement resistance decreased. Since the Si content of Production No. R was high, the precipitation of alloy carbides in the hot rolling process was suppressed, and the segregation of grain boundary strengthening elements to the prior austenite grain boundaries in the annealing process was suppressed, resulting in a decrease in hydrogen embrittlement resistance. Since the Mn content of Production No. S was low, the tensile strength was less than 1500 MPa. Production No. T had a high Mn content and E GB decreased, resulting in a decrease in hydrogen embrittlement resistance. Since the P content of Production No. U was high, E GB decreased and the hydrogen embrittlement resistance decreased. Since the S content of Production No. V was high, the hydrogen embrittlement resistance decreased. Since the N content of Production No. W was high, coarse nitrides were formed, resulting in a decrease in hydrogen embrittlement resistance. Since the Al content of Production No. X was high, coarse Al oxides were formed, resulting in a decrease in hydrogen embrittlement resistance. Manufacturing No. Y had a high B content, resulting in the formation of coarse B oxides and a decrease in hydrogen embrittlement resistance. Manufacturing No. Z had a high Ti content, leading to the formation of coarse carbonitrides and a decrease in hydrogen embrittlement resistance. Also, due to the formation of coarse carbonitrides, the amount of Ti segregating at the grain boundaries decreased, and E GB decreased. In addition, since the amount of C effective for improving strength decreased, the tensile strength was 1500 MPa or less. Manufacturing No. AA had a high Nb content, resulting in the formation of coarse Nb carbides and a decrease in hydrogen embrittlement resistance. Manufacturing No. AB had a high V content, leading to the formation of coarse V carbides and a decrease in hydrogen embrittlement resistance. Manufacturing No. AC had a high O content, resulting in the formation of oxides and a decrease in hydrogen embrittlement resistance. Manufacturing No. AD had a high Mo content, resulting in an increase in the precipitation of carbonitrides and a decrease in hydrogen embrittlement resistance. Also, due to the formation of coarse carbonitrides, the amount of Mo segregating at the grain boundaries decreased, and E GB decreased. Manufacturing No. AE had a high Cr content, resulting in the formation of coarse Cr carbides at the center segregation position in the steel and a decrease in hydrogen embrittlement resistance. Also, due to the formation of coarse carbonitrides, the amount of Cr segregating at the grain boundaries decreased, and E GB decreased. Manufacturing No. AF had a high Co content, resulting in the formation of coarse Co carbides and a decrease in hydrogen embrittlement resistance. Also, due to the formation of coarse carbonitrides, the amount of Co segregating at the grain boundaries decreased, and E GB decreased. Manufacturing No. AG had a high Ni content, resulting in a decrease in hydrogen embrittlement resistance. Manufacturing No. AH had a high Cu content, and E GB decreased to less than 0.50, resulting in a decrease in hydrogen embrittlement resistance. Manufacturing No. AI had a high W content, resulting in the formation of coarse W precipitates and a decrease in hydrogen embrittlement resistance. Also, the amount of W effective for grain boundary strengthening decreased, and E GB decreased to less than 0.50, resulting in a decrease in hydrogen embrittlement resistance. Since the Ta content in Production No. AJ was high, a large number of fine Ta carbides precipitated, resulting in a decrease in hydrogen embrittlement resistance. Since the Sn content in Production No. AK was high, E decreased below 0.50 due to grain boundary embrittlement, GB resulting in a decrease in hydrogen embrittlement resistance. Since the Sb and As contents in Production Nos. AL and AM were high respectively, E decreased below 0.50 due to grain boundary segregation, GB resulting in a decrease in hydrogen embrittlement resistance. Since the Mg and Ca contents in Production Nos. AN and AO were high respectively, the hydrogen embrittlement resistance decreased due to the formation of coarse inclusions. Since the Y, Zr, La, and Ce contents in Production Nos. AP - AS were high respectively, coarse oxides were formed, resulting in a decrease in hydrogen embrittlement resistance. In contrast, in Production Nos. A - O, by appropriately controlling the chemical composition and structure of the steel sheet and the grain boundary strength E GB of the prior austenite grain boundaries, a steel sheet with high strength and excellent hydrogen embrittlement resistance could be obtained.

[0093] <Example 2> Furthermore, in order to examine the influence of manufacturing conditions, for the steel grades (Steel Nos. A to O) in which excellent properties were observed in Example 1, in the same equipment as in Example 1, steel slabs were inserted into a furnace heated to 1250 to 1100 °C, held for 60 minutes, then taken out into the atmosphere, and hot-rolled steel sheets with a thickness of 2.3 mm were produced under the manufacturing conditions described in Tables 3-1 to 3-2. Furthermore, cold-rolled steel sheets were obtained with the conditions after coiling being as described in Tables 3-1 to 3-4. A part of the cold-rolled steel sheets was made into plated steel sheets with a plating layer formed. Here, the symbols GI and GA for the plating treatment indicate the method of zinc plating treatment. GI is a steel sheet in which the steel sheet was immersed in a molten zinc plating bath at 460 °C to form a zinc plating layer on the surface of the steel sheet, and GA is a steel sheet in which after the steel sheet was immersed in the molten zinc plating bath, the temperature of the steel sheet was raised to 485 °C to form an alloy layer of iron and zinc on the surface of the steel sheet. Plating was carried out as it was (without cooling to room temperature once) after the intermediate holding in the second temperature range. When the intermediate holding was not carried out in the second temperature range, it was carried out during the cooling from 25 to 300 °C. Also, in Tables 3-3 to 3-4, the examples described with "−" for tempering are examples without tempering. Also, the inter-pass time in the table is the inter-pass time for each pass with a reduction ratio of 20% or more (since rolling was carried out on a tandem rolling mill, each inter-pass time was the same). Also, in Tables 3-3 and 3-4, the holding time in the annealing post-cooling process is the holding time in the second temperature range when cooling was carried out to the second temperature range, but when the cooling stop temperature is outside the second temperature range, it is the holding time near that temperature.

[0094] For the obtained cold-rolled steel sheets (including plated steel sheets), in the same manner as in Example 1, the area ratios of ferrite, martensite, and tempered martensite, and the balance (one or more of bainite, pearlite, and retained austenite) in the microstructure were determined, and the concentrations of each alloy element on the austenite grain boundaries were measured, and E GB was determined.

[0095] Also, the tensile strength and total elongation of the obtained cold-rolled steel sheets were evaluated in the same manner as in Example 1.

[0096] Also, the hydrogen embrittlement resistance was evaluated by the following method. (Method for evaluating hydrogen embrittlement resistance) For the hot-dip galvanized steel sheet manufactured using the manufacturing method of the steel sheet according to the embodiment of the present invention, the hydrogen embrittlement resistance was evaluated by the following method. Specifically, after the steel sheet was sheared with a clearance of 15%, a U-bending test was performed at 8R. A strain gauge was attached to the center of the obtained test piece, and stress was applied by tightening both ends of the test piece with bolts. The applied stress was calculated from the strain of the monitored strain gauge. The load stress was set to a stress corresponding to 80% of the tensile strength (TS) (for example, in the case of A-1 in Table 4, the applied stress = 1540 MPa × 0.8 = 1232 MPa). This is because it is considered that the residual stress introduced during forming corresponds to the tensile strength of the steel sheet. The obtained U-bending test piece was immersed in an HCl aqueous solution with a liquid temperature of 25°C and a pH of 2, and held for 96 hours to check for cracks. Since the lower the pH of the HCl aqueous solution and the longer the immersion time, the larger the amount of hydrogen that penetrates into the steel sheet, the hydrogen embrittlement environment becomes severe. After immersion, the total crack length of the U-bending test piece was measured (when multiple cracks were observed, it was the sum of the individually measured values). The smaller the total crack length, the better the hydrogen embrittlement resistance. In particular, when a crack with a length exceeding 1.00 mm was observed, it was regarded as NG; when no crack was observed and when a minor crack with a length of 1.00 mm or less was observed, it was regarded as OK; among the OK cases, when no crack was observed and when the crack length was 0.70 mm or less, it was evaluated as Ex. The cases evaluated as OK and Ex were regarded as qualified, and the case of NG was regarded as unqualified. The results are shown in Table 4.

[0097]

Table 3-1

[0098]

Table 3-2

[0099]

Table 3-3

[0100]

Table 3-4

[0101]

Table 4

[0102] Referring to Table 4, for Production No. A-2, the time from the end of finish rolling to the start of cooling was long. Therefore, the ferrite transformation during the cooling process after finish rolling was suppressed, leading to coarsening of the pearlite structure and delay in the precipitation of alloy carbides. As a result, alloying elements contributing to the improvement of grain boundary strength could not fully segregate at the grain boundaries, and E GB fell below 0.50. Consequently, the hydrogen embrittlement resistance decreased. For Production No. B-2, since the rolling start temperature of hot rolling was low, unrecrystallized austenite remained, resulting in delay in the precipitation of alloy carbides, and E GB fell below 0.50. As a result, the hydrogen embrittlement resistance decreased. For Production No. C-2, since the rolling start temperature of hot rolling was high, the grain size of recrystallized austenite coarsened, leading to delay in the precipitation of alloy carbides, and E GB fell below 0.50. Consequently, the hydrogen embrittlement resistance decreased. For Production No. D-2, since the rolling end temperature was low, unrecrystallized austenite remained, causing delay in the precipitation of alloy carbides, and E GB fell below 0.50. As a result, the hydrogen embrittlement resistance decreased. For Production No. E-2, since the rolling end temperature was high, the grain size of recrystallized austenite coarsened, resulting in delay in the precipitation of alloy carbides, and E GB fell below 0.50. Consequently, the hydrogen embrittlement resistance decreased. For Production No. F-2, since the coiling temperature was low, the diffusion of alloying elements was delayed, and the precipitation of alloy carbides was suppressed, so E GBIt was less than 0.50. As a result, the hydrogen embrittlement resistance decreased. In Manufacturing No. G-2, since the coiling temperature was high, an internal oxide layer was formed on the surface of the hot-rolled steel sheet, and cracks occurred on the steel sheet surface during subsequent processing. Therefore, the tissue analysis and the evaluation of mechanical properties were not performed. In Manufacturing No. H-2, since the average cooling rate up to the coiling temperature was low, ferrite and pearlite transformations occurred and the precipitation of alloy carbides was suppressed, and E GB It was less than 0.50. As a result, the hydrogen embrittlement resistance decreased. In Manufacturing No. K-2, since the inter-pass time of the passes with a rolling reduction of 20% or more in finish rolling was long, the precipitation of alloy carbides in the hot-rolling process was delayed. As a result, E GB was less than 0.50, and the hydrogen embrittlement resistance decreased. In Manufacturing No. L-2, the holding temperature after the hot-rolling process was low, the residence time at 400 to 550 °C was less than 600 seconds, and the precipitation of alloy carbides did not occur sufficiently, so that E GB was less than 0.50. As a result, the hydrogen embrittlement resistance deteriorated. In Manufacturing No. M-2, the holding temperature after the hot-rolling process was high, the residence time at 400 to 550 °C was less than 600 seconds, and the precipitation of alloy carbides did not occur sufficiently, so that E GB was less than 0.50. As a result, the hydrogen embrittlement resistance decreased.

[0103] In Manufacturing No. A-3, since the annealing temperature was low, ferrite transformation progressed during holding, and the tensile strength was less than 1500 MPa. In Manufacturing No. B-3, since the annealing temperature was high, the concentration of alloying elements segregating at grain boundaries decreased. As a result, E GB was less than 0.50, and the hydrogen embrittlement resistance decreased. In Manufacturing No. C-3, since the average cooling rate from the annealing temperature was low, ferrite transformation occurred during cooling, and the tensile strength did not reach 1500 MPa. In Manufacturing No. F-3, since the cooling stop temperature was high, bainite transformation occurred, and the tensile strength did not reach 1500 MPa. Manufacturing No. H-3 had a high tempering temperature, resulting in softening of martensite and the tensile strength not reaching 1500 MPa. Manufacturing No. J-3 had a long tempering time, resulting in excessive softening of martensite and the tensile strength not reaching 1500 MPa. Manufacturing No. K-3 had low cooling stop temperature and intermediate holding temperature in the post-annealing cooling process (outside the second temperature range), resulting in bainite transformation and the tensile strength not reaching 1500 MPa. Manufacturing No. L-3 had high cooling stop temperature and intermediate holding temperature in the post-annealing cooling process (outside the second temperature range), resulting in ferrite transformation and pearlite transformation and the tensile strength not reaching 1500 MPa. Manufacturing No. M-3 had a long holding time at the second cooling stop temperature, resulting in bainite transformation during holding and the tensile strength not reaching 1500 MPa. Manufacturing No. N-3 had a low cooling rate from the second cooling stop temperature, resulting in ferrite transformation and bainite transformation during cooling and the tensile strength not reaching 1500 MPa.

[0104] In contrast, in all examples according to the present invention, by appropriately controlling hot rolling, coiling, annealing, etc., a steel sheet with high strength and excellent hydrogen embrittlement resistance could be obtained.

[0105] Figure 1 is a diagram showing the relationship between E GB and the tensile strength on the hydrogen embrittlement resistance of the steel sheets in Example 1 and Example 2. The ▲ in Figure 1 is an example where the hydrogen embrittlement resistance did not achieve the target, and the 〇 is an example where the hydrogen embrittlement resistance achieved the target. As shown in Figure 1, by setting E GB to 0.50 or more, excellent hydrogen embrittlement resistance can be obtained even for high-strength materials of 1500 MPa or more.

Industrial Applicability

[0106] According to the present invention, it is possible to provide a high-strength steel sheet having excellent hydrogen embrittlement resistance characteristics. When this steel sheet is applied to steel sheets for automobiles or the like, it contributes to reducing the weight of the vehicle body and improving fuel efficiency.

Claims

1. by mass percentage, C: 0.150 - 0.400%, Si: 0.01 - 2.00%, Mn: 0.80 - 2.00%, P: 0.0001 - 0.0200%, S: 0.0001 - 0.0200%, Al: 0.001 - 1.000%, N: 0.0001 - 0.0200%, O: 0.0001 - 0.0200%, Co: 0 - 0.500%, Ni: 0 - 1.000%, Mo: 0 - 1.000%, Cr: 0 - 2.000%, Ti: 0 - 0.500%, B: 0 - 0.0100%, Nb: 0 - 0.500%, V: 0 - 0.500%, Cu: 0 - 0.500%, W: 0 - 0.100%, Ta: 0 - 0.100%, Mg: 0 - 0.050%, Ca: 0 - 0.050%, Y: 0 - 0.050%, Zr: 0 - 0.050%, La: 0 - 0.050%, Ce: 0 - 0.050%, Sn: 0 - 0.050%, Sb: 0 - 0.050%, As: 0 - 0.050%, and the balance: Fe and impurities having a chemical composition consisting of wherein the microstructure, by 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 consisting of When the interface with an orientation difference of 15 deg. or more between adjacent martensite and tempered martensite is taken as the prior austenite grain boundary, the bonding strength energy E determined by the concentration of each alloy element on the prior austenite grain boundary ＧＢ satisfies the following formula (1): with a tensile strength of 1500 MPa or more, characterized in that it is a steel sheet. E ＧＢ = 1 + (3×[Co] + 0.7×[Ni] + 5.5×[Mo] + 0.7×[Cr] + 2.9×[Ti] + 47×[B] + 4.3×[Nb] + 4.5×[V] + 5.2×[W] + 3.1×[Ta] + 4.3×[Zr] - 0.25×[Mn] - 0.1×[P] - [Cu] - 1.1×[Sn] - 0.6×[Sb] - 0.9×[As]) ≥ 0.50 (1) Here, [chemical symbol] in the formula represents the concentration of each alloy element in mass% on the prior austenite grain boundary.

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

3. having a coating layer containing zinc, aluminum, magnesium or their alloys on the surface, The steel sheet according to claim 1 or 2, characterized in that...

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