JP7705069B2 - steel plate - Google Patents

Description

The present invention relates to a steel sheet.
This application claims priority based on Japanese Patent Application No. 2021-172424, filed on October 21, 2021, the contents of which are incorporated herein by reference.

Today, with the highly specialized fields of industrial technology, materials used in each field are required to have special and advanced performance. In particular, with regard to steel sheets for automobiles, the demand for high-strength steel sheets has increased significantly in order to reduce the weight of car bodies and improve fuel efficiency, due to consideration for the global environment. However, many metal materials have various properties that deteriorate as they are strengthened, and their susceptibility to hydrogen embrittlement in particular increases. It is known that steel members are particularly susceptible to hydrogen embrittlement when their tensile strength is 1200 MPa or more, and there have been cases of hydrogen embrittlement cracking in bolt steel, which has been strengthened ahead of the automotive field. Therefore, there is a strong demand for a fundamental solution to hydrogen embrittlement in high-strength steel sheets with a tensile strength of 1500 MPa or more.

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

In Patent Document 1, an ultra-high strength thin steel sheet with excellent hydrogen embrittlement resistance and workability is described that satisfies, in mass%, C: over 0.25 to 0.60%, Si: 1.0 to 3.0%, Mn: 1.0 to 3.5%, P: 0.15% or less, S: 0.02% or less, Al: 1.5% or less (excluding 0%), Mo: 1.0% or less (excluding 0%), Nb: 0.1% or less (excluding 0%), with the balance being iron and unavoidable impurities, and is processed at a tensile strength of 3%. The present invention discloses an ultra-high strength thin steel sheet having excellent hydrogen embrittlement resistance, characterized in that the metal structure after processing satisfies, in terms of area ratio relative to the total structure, retained austenite structure: 1% or more, bainitic ferrite and martensite: 80% or more in total, and ferrite and pearlite: 9% or less in total (including 0%), while also satisfying an average axial ratio (major axis/minor axis) of the retained austenite crystal grains: 5 or more, and having a tensile strength of 1,180 MPa or more.

Patent Document 2 discloses a high-strength steel plate having a tensile strength of 1,500 MPa or more, the steel components containing Si + Mn: 1.0% or more, the main phase structure being layers of ferrite and carbides, the aspect ratio of the carbides being 10 or more and the layer spacing being 50 nm or less being 65% or more in volume ratio relative to the entire structure, and further, the fraction of carbides forming layers with ferrite that have an aspect ratio of 10 or more and have an angle of 25° or less with respect to the rolling direction being 75% or more in area ratio, thereby providing a high-strength steel plate with excellent bendability in the rolling direction and resistance to delayed fracture.

Patent Document 3 discloses, as a thin ultra-high strength cold rolled steel sheet excellent in bendability and delayed fracture resistance, an ultra-high strength cold rolled steel sheet excellent in bendability, characterized in that 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 unavoidable impurities, has a steel sheet surface soft portion that satisfies the relationship "hardness of steel sheet surface soft portion/hardness of steel sheet center portion≦0.8", the proportion of the steel sheet surface soft portion in the sheet thickness is 0.10 to 0.30, the volume fraction of the steel sheet surface soft portion is tempered martensite 90% or more, the structure of the steel sheet center portion is tempered martensite, and has a tensile strength of 1270 MPa or more.

Patent Document 4 discloses a cold-rolled steel sheet having a tensile strength of 1,470 MPa or more and excellent bending workability and delayed fracture resistance, the cold-rolled steel sheet containing, by mass%, C: 0.15-0.20%, Si: 1.0-2.0%, Mn: 1.5-2.5%, P: 0.020% or less, S: 0.005% or less, Al: 0.01-0.05%, N: 0.005% or less, Ti: 0.1% or less, Nb: 0.1% or less, B: 5-30 ppm, with the balance being Fe and unavoidable impurities, and having a metal structure in which the volume fraction of the tempered martensite phase is 97% or more and the volume fraction of the retained austenite phase is less than 3%.

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

Patent Document 6 describes an ultra-high strength cold rolled steel sheet having excellent hydrogen embrittlement resistance and a tensile strength of 1300 MPa or more, which contains C: 0.150 to 0.300%, Si: 0.001 to 2.0%, Mn: 2.10 to 4.0%, P: 0.05% or less, S: 0.01% or less, N: 0.01% or less, Al: 0.001% to 1.0%, Ti: 0.001% to 0.10%, B: 0.0001% to 0.20%, and H: 0.0001% to 0.05%. The present invention discloses an ultra-high strength cold rolled steel sheet having a steel structure containing 0.010% of iron, in which the amount of solute B solB [mass %] and the prior austenite grain size Dγ [μm] satisfy the relationship solB·Dγ≧0.0010, the amount of polygonal ferrite is 10% or less, the amount of bainite is 30% or less, the amount of retained austenite is 6% or less, the amount of tempered martensite is 60% or more, the number density of Fe carbides in the tempered martensite is 1×10 ⁶ /mm ² or more, the average dislocation density of the entire steel is 1.0×10 ¹⁵ to 2.0×10 ¹⁶ /m ² , and the grain size is 7.0 μm or less.

Japanese Patent Application Publication No. 2006-207019 Japanese Patent Application Publication No. 2010-138489 Japanese Patent Application Publication No. 2011-179030 Japanese Patent Application Publication No. 2010-215958 Japanese Patent Application Publication No. 2016-153524 Japanese Patent Application Publication No. 2016-050343

As described above, several techniques have been proposed to improve the hydrogen embrittlement resistance of high-strength steel sheets. However, Patent Document 1 only discloses the hydrogen embrittlement resistance when a stress of 1000 MPa is applied, and does not provide any technical solution for the hydrogen embrittlement resistance when a higher stress is applied.
As described above, hydrogen embrittlement occurs when hydrogen accumulates at grain boundaries, reducing the bond strength of the grain boundaries. Therefore, it is believed that if the bond strength of the grain boundaries can be increased, cracks due to hydrogen embrittlement can be suppressed. However, Patent Documents 1 to 6 do not consider a method for improving hydrogen embrittlement resistance from this perspective. In recent years, requirements for hydrogen embrittlement resistance have become even stricter, and Patent Documents 1 to 6 may not be able to meet such strict requirements.
That is, conventionally, there has been room for improvement in hydrogen embrittlement resistance in high-strength steel plates having a microstructure mainly composed of martensite and tempered martensite.

Furthermore, in Patent Document 2, a steel sheet has a pearlite structure as the main phase, a ferrite phase in the remaining structure is 20% or less in volume ratio relative to the entire structure, and the lamellar spacing of the pearlite structure is 500 nm or less, and the steel sheet has a Vickers hardness of HV200 or more, and is obtained by cold rolling at a rolling reduction 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.
Furthermore, in Patent Document 3, in order to improve delayed fracture properties, it is necessary to hold the material at 650° C. or 700° C. for 20 minutes or more in an atmosphere with a dew point of 15° C. or higher, which also results in a problem of low productivity.

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

As described above, hydrogen embrittlement is considered to be a crack that occurs from the grain boundary due to the segregation of hydrogen in steel to the grain boundary, which reduces the bond strength of the grain boundary. Therefore, the present inventors have focused on the bond strength of the grain boundary and conducted various studies on a method for improving hydrogen embrittlement resistance.
As a result, the inventors discovered that by segregating a specific alloy element to the grain boundary, the grain boundary bonding strength is improved and infiltrated hydrogen is less likely to segregate to the grain boundary, and even if hydrogen does invade, it is possible to suppress a decrease in the grain boundary bonding strength due to hydrogen.

The present invention has been made in view of the above findings.
[1] A steel sheet according to one embodiment 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. 0001-0.0200%, 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, and the microstructure is composed of, in terms of area ratio, ferrite: 5.0% or less, martensite and tempered martensite: a total of more than 90.0%, and the balance: one or more of bainite, pearlite, and retained austenite, and the orientation difference between adjacent martensite and tempered martensite is 15 deg. When the above interface is regarded as a prior austenite grain boundary, the bond strength energy E _GB determined by the concentration of each alloy element on the prior austenite grain boundary satisfies the following formula (1), and the tensile strength is 1500 MPa or more.
_EGB =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, the [chemical symbols] in the formula represent the concentrations of the respective alloying elements in mass % on the prior austenite grain boundaries.
[2] The steel sheet according to [1], wherein the chemical composition is Co: 0.01 to 0.500%, Ni: 0.01 to 1.000%, Mo: 0.01 to 1.000%, Cr: 0.001 to 2.000%, Ti: 0.001 to 0.500%, B: 0.0001 to 0.0100%, Nb: 0.001 to 0.500%, V: 0.001 to 0.500%, Cu: 0.001 to 0.500%, W: 0.001 to 0.100%, Ta: 0.0 It may contain one or more selected from the group consisting of: 0.01 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%.
[3] The steel sheet according to [1] or [2] may have a coating layer containing zinc, aluminum, magnesium or an alloy thereof on a surface thereof.

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

FIG. 1 is a diagram showing the relationship between hydrogen embrittlement resistance, _EGB and tensile strength of steel plates in examples of the present invention.

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

<Chemical composition>
First, the range of the content of each element constituting the chemical composition of the steel sheet according to this embodiment will be described. Hereinafter, "%" relating to the content of an element means "mass%". In addition, the range indicated with "~" includes both ends as the lower limit or upper limit.

C: 0.150-0.400%
C is an effective element for increasing tensile strength at low cost. If the C content is less than 0.150%, the target tensile strength cannot be obtained, and the fatigue properties of the welded portion deteriorate. For this reason, the C content is set to 0.150% or more. The C content may be 0.160% or more, 0.180% or more, or 0.200% or more.
On the other hand, if the C content exceeds 0.400%, hydrogen embrittlement resistance and weldability are reduced. 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.

Si: 0.01~2.00%
Silicon acts as a deoxidizer and is an element that affects the morphology of carbides and residual austenite after heat treatment. If the silicon content is less than 0.01%, it becomes difficult to suppress the generation of coarse oxides. These coarse oxides become the starting points of cracks, and the hydrogen embrittlement resistance deteriorates as the cracks propagate through the steel material. For this reason, the silicon content is set to 0.01% or more. The silicon content may be 0.05% or more, 0.10% or more, or 0.30% or more.
On the other hand, if the Si content exceeds 2.00%, 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.

Mn: 0.80-2.00%
Mn is an element effective in increasing the strength of a steel sheet. If the Mn content is less than 0.80%, the effect is not sufficiently obtained. 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, if the Mn content exceeds 2.00%, Mn not only promotes co-segregation with P and S, but also may deteriorate the corrosion resistance and hydrogen embrittlement resistance. 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.

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

S: 0.0001-0.0200%
S is an element that generates nonmetallic inclusions such as MnS in steel. If the S content exceeds 0.0200%, the generation of nonmetallic inclusions that become the starting point of cracks during cold working becomes significant. In this case, even if the grain boundaries are strengthened, cracks will occur from the nonmetallic inclusions, and these cracks will propagate through the steel material, deteriorating the hydrogen embrittlement resistance. For this reason, 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 better. However, if the S content is less than 0.0001%, the time required for refining increases, leading to a significant increase in costs. For this reason, the S content is set to 0.0001% or more. The S content may be 0.0005% or more, 0.0010% or more, or 0.0020% or more.

Al: 0.001-1.000%
Al acts as a deoxidizer for steel and is an element that stabilizes ferrite. If the Al content is less than 0.001%, the effect is not sufficiently obtained. Therefore, the Al content is set to 0.001% or more. The Al content may be 0.005% or more, 0.010% or more, or 0.020% or more.
On the other hand, when the Al content exceeds 1.000%, coarse Al oxides are generated. These coarse oxides become the starting points of cracks. Therefore, when coarse Al oxides are generated, even if the grain boundaries are strengthened, cracks occur in the coarse oxides, and these cracks propagate through the steel material, deteriorating the hydrogen embrittlement resistance. For this reason, the Al content is set to 1.000% or less. The Al content may be 0.950% or less, 0.900% or less, or 0.800% or less.

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

O: 0.0001-0.0200%
O is an element that forms oxides and deteriorates hydrogen embrittlement resistance. In particular, oxides often exist as inclusions, and when present on a punched end surface or a cut surface, they form notched scratches or coarse dimples on the end surface, which leads to stress concentration during strong processing and becomes the starting point for crack formation, resulting in significant deterioration of workability. If the O content exceeds 0.0200%, the tendency of the above-mentioned deterioration of workability becomes significant. For this reason, the O content is set to 0.0200% or less. The O content may be 0.0180% or less, 0.0150% or less, or 0.0100% or less.
A small O content is preferable. However, an O content of less than 0.0001% leads to excessive cost increase and is economically undesirable. For this reason, the O content is set to 0.0001% or more. The O content may be 0.0005% or more, 0.0010% or more, or 0.0015% or more.

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

Co: 0-0.500%
Co is an element effective for controlling the morphology of carbides and increasing the strength of the steel sheet. Co is also an element that contributes to improving the bonding strength of grain boundaries. Therefore, Co may be contained. To obtain a sufficient effect, the Co content is preferably 0.010% or more. The Co content may be 0.020% or more, 0.050% or more, or 0.100% or more.
On the other hand, if the Co content exceeds 0.500%, coarse Co carbides are precipitated. In this case, hydrogen embrittlement resistance 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.

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

Mo: 0~1.000%
Mo is an element effective in increasing the strength of steel sheet. Mo is also an element that has the effect of suppressing ferrite transformation that occurs during heat treatment in a continuous annealing facility or a continuous hot-dip galvanizing facility. Mo is also an element that contributes to improving the bonding strength of grain boundaries. Therefore, Mo may be contained. To obtain the above effect, the Mo content is preferably 0.010% or more. The Mo content may be 0.020% or more, 0.050% or more, or 0.080% or more.
On the other hand, if the Mo content exceeds 1.000%, the effect of suppressing the ferrite transformation is saturated. Therefore, the Mo content is set to 1.000% or less. The Mo content may be 0.900% or less, 0.800% or less, or 0.600% or less.

Cr: 0-2.000%
Cr, like Mn, is an element that suppresses pearlite transformation and is effective in increasing the strength of steel. Cr is also an element that contributes to improving the bonding strength of grain boundaries. Therefore, Cr may be contained. To obtain the above effect, the Cr content is preferably 0.001% or more. The Cr content may be 0.005% or more, 0.010% or more, or 0.050% or more.
On the other hand, if the Cr content exceeds 2.000%, coarse Cr carbides are formed in the central segregation portion, and hydrogen embrittlement resistance may deteriorate. Therefore, 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.

Ti: 0~0.500%
Ti is an element that contributes to increasing the strength of the steel sheet through precipitation strengthening, fine grain strengthening by inhibiting the growth of ferrite crystal grains, and dislocation strengthening through inhibiting recrystallization. Ti is also an element that contributes to improving the bonding strength of grain boundaries. Therefore, Ti may be contained. To obtain the above effect, the Ti content is preferably 0.001% or more. The Ti content may be 0.003% or more, 0.010% or more, or 0.050% or more.
On the other hand, if the Ti content exceeds 0.500%, the precipitation of carbonitrides increases and hydrogen embrittlement resistance may deteriorate. Therefore, the Ti content is set to 0.500% or less. The Ti content may be 0.450% or less, 0.400% or less, or 0.300% or less.

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

Nb: 0-0.500%
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. Nb is also an element that contributes to improving the bonding strength of grain boundaries. Therefore, Nb may be contained. To obtain the above effects, the Nb content is preferably 0.001% or more. The Nb content may be 0.002% or more, 0.010% or more, or 0.020% or more.
On the other hand, if the Nb content exceeds 0.500%, the generation of coarse Nb carbides becomes significant. Since the coarse Nb carbides are prone to cracking, the generation of the coarse Nb carbides may deteriorate the hydrogen embrittlement resistance. For this reason, 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.

V: 0-0.500%
V is an element that contributes to increasing the strength of the steel sheet through precipitation strengthening, fine grain strengthening by inhibiting the growth of ferrite crystal grains, and dislocation strengthening through inhibiting recrystallization. V also contributes to improving the bonding strength of grain boundaries. Therefore, V may be contained. To obtain the above effect, the V content is preferably 0.001% or more. The V content may be 0.002% or more, 0.010% or more, or 0.020% or more.
On the other hand, if the V content exceeds 0.500%, the precipitation of carbonitrides increases and hydrogen embrittlement resistance may deteriorate. Therefore, the V content is set to 0.500% or less. The V content may be 0.450% or less, 0.400% or less, or 0.300% or less.

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

W: 0~0.100%
W is an element effective in increasing the strength of steel sheet. In addition, W forms precipitates and crystallized products. Precipitates and crystallized products containing W become hydrogen trapping sites, so W is an element effective in improving hydrogen embrittlement resistance. In addition, W is an element that contributes to improving the bonding strength of grain boundaries. Therefore, W may be contained. To obtain the above effects, the W content is preferably 0.001% or more. The W content may be 0.002% or more, 0.005% or more, or 0.010% or more.
On the other hand, if the W content exceeds 0.100%, the generation of coarse W precipitates or crystallized particles becomes significant. These coarse W precipitates or crystallized particles are prone to cracking, and these cracks propagate through the steel material under low load stress. Therefore, if coarse W precipitates or crystallized particles are generated, hydrogen embrittlement resistance 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.

Ta: 0~0.100%
Ta, like Nb, V, and W, is an element that is effective in controlling the morphology of carbides and increasing the strength of the steel sheet. Ta is also an element that contributes to improving the bonding strength of grain boundaries. Therefore, Ta may be contained. To obtain the above effects, the Ta content is preferably 0.001% or more. The Ta content may be 0.002% or more, 0.005% or more, or 0.010% or more.
On the other hand, if the Ta content exceeds 0.100%, a large number of fine Ta carbides are precipitated, and as the strength of the steel sheet increases, the ductility and the bending resistance and hydrogen embrittlement resistance 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.

Mg: 0-0.050%
Mg is an element that can control the morphology of sulfides with a small amount of content. Therefore, Mg may be contained. To obtain the above effects, the Mg content is preferably 0.001% or more. The Mg content may be 0.005% or more, 0.010% or more, or 0.020% or more.
On the other hand, if the Mg content exceeds 0.050%, coarse inclusions may be formed and hydrogen embrittlement resistance may be deteriorated. Therefore, 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.

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

Y: 0~0.050%
Y is an element that can control the form of sulfides by containing a small amount of Y, similar to Mg and Ca. Therefore, Y may be contained. To obtain the above effects, the Y content is preferably 0.001% or more. The Y content may be 0.002% or more, 0.004% or more, or 0.006% or more.
On the other hand, if the Y content exceeds 0.050%, coarse Y oxides are generated, and hydrogen embrittlement resistance may deteriorate. Therefore, the Y content is set to 0.050% or less. The Y content may be 0.040% or less, 0.030% or less, or 0.020% or less.

Zr: 0~0.050%
Zr is an element that can control the morphology of sulfides by containing a small amount of it, similar to Mg, Ca, and Y. Zr is also an element that contributes to improving the bonding strength of grain boundaries. Therefore, Zr may be contained. To obtain the above effects, the Zr content is preferably 0.001% or more. The Zr content may be 0.002% or more, 0.004% or more, or 0.006% or more.
On the other hand, if the Zr content exceeds 0.050%, coarse Zr oxides are generated, and hydrogen embrittlement resistance may deteriorate. Therefore, 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.

La: 0-0.050%
La is an element that can control the form of sulfides by containing a small amount of it, similar to Mg, Ca, Y, and Zr. Therefore, La may be contained. To obtain the above effects, the La content is preferably 0.001% or more. The La content may be 0.002% or more, 0.004% or more, or 0.006% or more.
On the other hand, if the La content exceeds 0.050%, La oxides are generated and hydrogen embrittlement resistance may deteriorate. 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.

Ce: 0~0.050%
Ce is an element that can control the form of sulfides by containing a small amount of Ce, similar to La. Therefore, Ce may be contained. To obtain the above effects, the Ce content is preferably 0.001% or more. The Ce content may be 0.002% or more, 0.004% or more, or 0.006% or more.
On the other hand, if the Ce content exceeds 0.050%, Ce oxides are generated and hydrogen embrittlement resistance may deteriorate. 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.

Sn: 0-0.050%
Sn is an element contained in steel when scrap is used as a raw material. If the Sn content is high, the hydrogen embrittlement resistance may be deteriorated due to embrittlement of grain boundaries. This adverse effect is particularly noticeable when the Sn content exceeds 0.050%. For this reason, 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 better, and 0% may be used. However, if the Sn content is less than 0.001%, the refining cost increases. Therefore, the Sn content may be 0.001% or more. The Sn content may be 0.002% or more, 0.005% or more, or 0.010% or more.

Sb: 0 to 0.050%
Sb is an element contained when scrap is used as a steel raw material, similar to Sn. Sb is an element that strongly segregates at grain boundaries, leading to embrittlement of the grain boundaries and a decrease in ductility. This adverse effect is particularly significant when the Sb content exceeds 0.050%. For this reason, 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 better, and it may be 0%, but if the Sb content is less than 0.001%, the refining cost increases. Therefore, the Sb content may be 0.001% or more. The Sb content may be 0.002% or more, 0.005% or more, or 0.008% or more.

As: 0~0.050%
Like Sn and Sb, As is an element contained when scrap is used as a steel raw material, and strongly segregates at grain boundaries, leading to embrittlement of the grain boundaries and a decrease in ductility. If the As content is high, hydrogen embrittlement resistance may decrease. This adverse effect is particularly noticeable when the As content exceeds 0.050%. For this reason, 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 As content is preferably as low as possible, and may be 0%, but if the As content is less than 0.001%, refining costs increase. 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.

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

The chemical composition of the steel plate 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 the chips in accordance with JIS G1201:2014. In this case, the chemical composition is the average content in the entire plate thickness. C and S, which cannot be measured by ICP-AES, may be measured using the combustion-infrared absorption method, N may be measured using the inert gas fusion-thermal conductivity method, and O may be measured using the inert gas fusion-non-dispersive infrared absorption method.
When the steel sheet has a coating layer on its surface, the coating layer may be removed by mechanical grinding or the like before the chemical composition is analyzed. When the coating layer is a plating layer, the plating layer may be removed by dissolving it in an acid solution to which an inhibitor that suppresses corrosion of the steel sheet has been added.

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

Ferrite: 5.0% or less Ferrite affects the deformability of steel mainly composed of martensite. As the area ratio of ferrite increases, local deformability and hydrogen embrittlement resistance decrease. In particular, when the area ratio of ferrite exceeds 5.0%, hydrogen embrittlement resistance may decrease due to fracture in elastic deformation under stress load. For this reason, 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%, a high level of control is required in the production and leads to a decrease in yield. Therefore, the area ratio of ferrite may be set to 1.0% or more.

Martensite and tempered martensite: total of more than 90.0% The total area ratio of martensite and tempered martensite affects the strength of the steel, and the larger 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 destruction during elastic deformation under stress load, or may cause deterioration of hydrogen embrittlement resistance due to the formation of structures other than martensite and tempered martensite and increased heterogeneity of the microstructure. For this reason, the total area ratio of martensite and tempered martensite is made to be 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%.

Remainder: one or more of bainite, pearlite, and retained austenite. The area ratio of the structure other than the above-mentioned structure (the remainder structure) may be 0%, but when a remainder structure exists, the remainder structure is one or more of bainite, pearlite, and retained austenite.
If the area ratio of the remaining structure exceeds 8.0%, it may lead to fracture in elastic deformation under stress load, and hydrogen embrittlement resistance may decrease. Therefore, the area ratio of the remaining structure is preferably 8.0% or less, and more preferably 7.0% or less. Among these, pearlite and retained austenite are particularly structures that deteriorate the local ductility of the steel, so the smaller the amount, the better.
On the other hand, in order to make the area ratio of the remaining structure 0%, a high level of control is required in the production, which may result in a decrease in yield. Therefore, the area ratio of the remaining structure may be 1.0% or more.

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

(Method of evaluating the area ratio of ferrite)
The area ratio of ferrite is obtained by observing the t/4 part (a range of 1/8 to 3/8 of the plate thickness centered at a position of 1/4 of the plate thickness from the surface in the plate thickness direction) by an electron channeling contrast image using a field emission scanning electron microscope (FE-SEM). The electron channeling contrast image is a method for detecting the crystal orientation difference within a crystal grain as a difference in image contrast, and in the image, a part that appears with a uniform contrast in a 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 eight fields of view of 35 μm×25 μm electron channeling contrast images is calculated by image analysis, and the average value is taken as the area ratio of ferrite.

(Method for evaluating the total area ratio of martensite and tempered martensite)
The total area ratio of martensite and tempered martensite is also obtained from the image taken by the above-mentioned electron channeling contrast. These structures are less easily etched than ferrite, so they are present as convex parts on the structure observation surface. Tempered martensite is a collection 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 multiple variants, that is, multiple iron-based carbide groups elongated in different directions. In addition, retained austenite is also present as convex parts on the structure observation surface. Therefore, by subtracting the area ratio of the convex parts obtained by the above procedure from the area ratio of the retained austenite measured by the procedure described later, it is possible to correctly measure the total area ratio of martensite and tempered martensite.

(Method for evaluating the total area ratio of bainite, pearlite and retained austenite)
The area ratio of the retained austenite can be calculated by measurement using X-rays. That is, the sample is removed from the plate surface to a position of 1/4 of the plate thickness in the plate thickness direction by mechanical polishing and chemical polishing. Then, the structure fraction of the retained austenite is calculated from the integrated intensity ratio of the diffraction peaks of (200), (211) of the bcc phase and (200), (220), and (311) of the fcc phase obtained by using MoKα rays as characteristic X-rays for the polished sample, and this is the area ratio of the retained austenite.
The area ratio of pearlite is calculated from the image taken by the above-mentioned electron channeling contrast. Pearlite is a structure in which plate-shaped carbides and ferrite are arranged.
Bainite is a collection of lath-shaped crystal grains that does not contain iron-based carbides with a major axis of 20 nm or more, or contains iron-based carbides with a major axis of 20 nm or more, and the carbides belong to a single variant, i.e., a group of iron-based carbides elongated in the same direction. Here, the group of iron-based carbides elongated in the same direction refers to iron-based carbides whose elongation directions differ by 5° or less.

<Bond Strength Energy>
Cracking due to hydrogen embrittlement occurs when hydrogen in steel segregates to grain boundaries, reducing the grain boundary bond strength, and the grain boundaries with reduced bond strength become the starting point of cracking. In contrast, segregating a specific alloying element to the grain boundaries improves the grain boundary bond strength. In addition, since the alloying elements are already segregated, hydrogen that has penetrated is less likely to segregate to the grain boundaries, and even if hydrogen penetrates, it is possible to suppress the decrease in grain boundary bond strength due to hydrogen.
The present inventors investigated the contribution of each alloying element to improving grain boundary strength in a microstructure mainly composed of martensite and tempered martensite, taking 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 degrees or more, as the main grain boundaries. As a result, they found that the bond strength energy E _GB can be expressed as the following formula (1) using the concentration of each alloying element on the grain boundary, and that when E _GB is 0.50 or more, hydrogen embrittlement resistance is clearly improved.
Therefore, in the steel sheet according to the present embodiment, when an interface between adjacent martensite and tempered martensite with an orientation difference of 15 degrees or more is defined as a prior austenite grain boundary, the bond strength energy E _GB determined by the concentration of each alloy element on the prior austenite grain boundary satisfies the following formula (1).
_EGB =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, the [chemical symbols] in the formula represent the concentration of each alloying element in mass % on the prior austenite grain boundary.
The reason why the interface with an orientation difference of 15 degrees or more was targeted is that hydrogen tends to accumulate preferentially at prior austenite grain boundaries with an orientation difference of 15 degrees or more.
As can be seen from formula (1), not all alloy elements that segregate to the grain boundaries increase the bond strength energy, but by segregating a large amount of alloy elements that increase the grain boundary bond energy, the grain boundary bond energy increases.
Here, the interface between adjacent martensite and tempered martensite, which has an orientation difference of 15 degrees or more, refers to an interface between martensite and martensite, which has an orientation difference of 15 degrees or more, an interface between martensite and tempered martensite, which has an orientation difference of 15 degrees or more, and an interface between tempered martensite and tempered martensite, which has an orientation difference of 15 degrees or more.

The concentration of each alloy element on the prior austenite grain boundary is determined by observing the t/4 portion (a range of 1/8 to 3/8 of the plate thickness centered at a position of 1/4 of the plate thickness from the surface in the plate thickness direction) using an EDS (energy dispersive X-ray spectroscope) of a TEM (transmission electron microscope) in the same manner as the above-mentioned SEM observation. More specifically, a spherical aberration corrected TEM transmission electron microscope (Cs-corrected TEM) is used as the TEM (transmission electron microscope).
Thin section samples used for TEM observation are obtained as follows: A sample for measuring the amount of alloying elements is taken from the steel plate sample in the range of 1/8 to 3/8 of the plate thickness, and polished to a thickness of about 100 μm by wet polishing using emery paper.
Then, the sample is electropolished by twin-jet electropolishing until it is thin enough for TEM observation. The electropolishing method is performed using a twin-jet electropolishing device. The appropriate conditions for twin-jet electropolishing vary depending on the base material components of the sample, so they must be extracted for each sample. After performing the twin jet, the thin sample is uniformly milled using Ar ion milling, improving the quantification accuracy of elements on the prior austenite grain boundaries.
The thus obtained thin sample 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 a bright-field image in TEM observation, when observation is performed at 30,000 times magnification, the prior austenite grain boundary, packet boundary, and block boundary appear as black lines. The sample is tilted and rotated so that the black line indicating any prior austenite grain boundary among these black lines is horizontal to the direction of incidence of the electron beam of the TEM. In this state, elemental analysis using EDS is performed at 100,000 times magnification directly above the prior austenite grain boundary. Elemental analysis using the number of integration times of EDS analysis is performed in the following manner. Point analysis is performed three 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 five prior austenite grain boundaries, and the average alloy element concentration is calculated. This average alloying element concentration is defined as the alloying element concentration on the prior austenite grain boundary.

(Mechanical properties)
In the steel plate according to this embodiment, the tensile strength (TS) is set to 1500 MPa or more as a strength that contributes to reducing the weight of an automobile body.
There is no need to set an upper limit, but since an increase in tensile strength may result in a decrease in formability, the tensile strength may be set to 2000 MPa or less.

(Thickness)
The steel plate according to the present embodiment is not limited in 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.

(Coating layer)
The steel sheet according to the present embodiment may have a coating layer containing zinc, aluminum, magnesium or an alloy thereof on one or both surfaces. This coating layer may consist of zinc, aluminum, magnesium or an alloy thereof and impurities.
By providing a coating layer on the surface, corrosion resistance is improved. In the case of automotive steel sheets, there are cases where the sheet thickness cannot be reduced below a certain level even if the sheet is made stronger due to concerns about holes caused by corrosion. Since one of the purposes of increasing the strength of steel sheets is to reduce weight by making them thinner, even if a high-strength steel sheet is developed, if the corrosion resistance is low, the locations where it can be used will be limited. One method that can be considered to solve these problems is to form a coating layer on the front and back surfaces to improve corrosion resistance.
Even if the coating layer is formed, the hydrogen embrittlement resistance of the steel plate according to this embodiment is not impaired.
The coating layer is, for example, a hot-dip galvanized layer, a galvannealed layer, an electrolytic galvanized layer, an aluminum plated layer, a Zn-Al alloy plated layer, an Al-Mg alloy plated layer, or a Zn-Al-Mg alloy plated layer.
When a coating layer is present on the surface, the reference surface of the t/4 portion described above is the surface of the base steel excluding the coating layer.

<Manufacturing method>
The steel sheet according to this embodiment can obtain the above-mentioned effects as long as it has the above-mentioned characteristics regardless of the manufacturing method, but it can be manufactured by a manufacturing method including the following steps (I) to (VII).
(I) a heating step of heating a steel slab having a predetermined chemical composition;
(II) a hot rolling step of hot rolling the heated steel slab to obtain a hot rolled steel sheet;
(III) A cooling step in which cooling of the hot-rolled steel sheet is started within 3.0 seconds from the completion of the hot rolling step, and the sheet 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;
(IV) a coiling step of coiling the hot-rolled steel sheet after the cooling step at the coiling temperature;
(V) a retention step of retaining the hot-rolled steel sheet after the coiling step in a temperature range of 400 to 550° C. for 600 seconds or more;
(VI) a cold rolling step of subjecting the hot-rolled steel sheet after the retention step to pickling and cold rolling to obtain a cold-rolled steel sheet;
(VII) An annealing step of annealing the cold-rolled steel sheet after the cold rolling step at an annealing temperature of 800° C. or more and less than 900° C.
Preferred conditions for each step will now be described.

(Heating process)
In the heating step, a steel billet such as a slab having the same chemical composition as the steel plate according to the present 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, 1000 to 1300°C.
From the viewpoint of productivity, the steel billets used are preferably cast by a continuous casting method, but may also be produced by an ingot casting method or a thin slab casting method.
When the steel slab obtained by continuous casting can be subjected to the hot rolling step while still at a sufficiently high temperature, the heating step may be omitted.

(Hot rolling process)
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, and in the finish rolling, multiple passes are performed, of which at least four passes are large reduction passes with a reduction rate of at least 20%, and the inter-pass time between the large reduction passes is 5.0 seconds or less. The rolling start temperature is 950 to 1100°C, and the rolling end temperature is 800 to 950°C.

[In the finish rolling, large reduction passes with a reduction rate of 20% or more: 4 passes or more]
[Inter-pass time: within 5.0 seconds]
By controlling the reduction ratio, the number of rolling passes, and the time between passes in the finish rolling, it is possible to control the morphology of the austenite grains to be equiaxed and fine. When the austenite grains become equiaxed and fine, the grain boundary diffusion of the alloy elements is promoted, and the precipitation of alloy carbides or nitrides is promoted at the grain boundaries. If the number of passes (large reduction passes) with a reduction ratio of 20% or more is less than four passes, unrecrystallized austenite remains, and therefore sufficient effect cannot be obtained. For this reason, the reduction ratio is set to 20% or more in four passes or more (four or more passes are performed with a reduction ratio of 20% or more). Preferably, the reduction ratio is set to 20% or more in five passes or more. On the other hand, the upper limit of the number of passes with a reduction ratio of 20% or more is not particularly limited, but in order to achieve more than 10 passes, it is necessary to install a large number of rolling stands, which may lead to an increase in the size of the equipment and an increase in manufacturing costs. For this reason, the number of passes (number of passes) with a reduction ratio of 20% or more may be 10 passes or less, 9 passes or less, or 7 passes or less.
In addition, the interpass time in the finish rolling has a large effect on the recrystallization and grain growth of austenite grains after rolling. Even when the number of large reduction passes is four or more, if the interpass time exceeds 5.0 seconds, grain growth is likely to occur and the austenite grains become coarse.
On the other hand, although it is not necessary to set a lower limit for the interpass time, if the interpass time for each heavy reduction pass is less than 0.2 seconds, the recrystallization of austenite may not be completed, the proportion of unrecrystallized austenite may increase, and sufficient effects may not be obtained. For this reason, it is preferable to set the interpass time for heavy reduction passes to 0.2 seconds or more. The interpass time may be 0.3 seconds or more or 0.5 seconds or more.

(cooling process)
In the cooling step, cooling of the hot-rolled steel sheet after the hot rolling step is started within 3.0 seconds from the completion of the hot rolling step (completion of the final pass of finish rolling), and 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.
If the time from the end 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./second, austenite transforms into ferrite before coiling. In this case, the driving force of the precipitates becomes small, and it becomes difficult for the 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 hardened phase is likely to form, which significantly deteriorates the subsequent manufacturability and reduces productivity.
In addition, if the cooling stop temperature is less than 550°C, the precipitation of precipitates is delayed, leading to deterioration of manufacturability and lowering productivity. In addition, if the cooling stop temperature is more than 700°C, austenite to ferrite transformation occurs, the driving force for carbide precipitation becomes smaller, and it becomes difficult for precipitates to precipitate uniformly and finely in the subsequent process, which is not preferable. In addition, if the cooling stop temperature is more than 700°C, an internal oxide layer is likely to be formed on the steel sheet surface, cracks are likely to occur on the surface, and the productivity in the subsequent pickling process is significantly deteriorated, which is not preferable.
There is no need to set a lower limit for the time from the end of hot rolling to the start of rolling, and the time may be as short as possible within the limits of the equipment.

(winding process)
(Staying process)
In the coiling process, the hot-rolled steel sheet after the cooling process is coiled at a coiling temperature (cooling stop temperature). In the subsequent retention process, the coiled hot-rolled steel sheet is held (retained) in a temperature range of 400 to 550°C for 600 seconds or more. By controlling the coiling and retention conditions, alloy carbides or nitrides are precipitated in the steel sheet.
The precipitates formed here can be made to be concentrated in the prior austenite grain boundaries by controlling the subsequent processes.
If the holding temperature is too high, the precipitates become coarse and are not uniformly dispersed. If the holding temperature is too low, the precipitates become fine, but it takes a long time to complete precipitation, which reduces manufacturability and productivity. If the holding time is too short, the alloy carbides do not precipitate sufficiently.
In order to keep the mixture under the above conditions, for example, a cover may be placed on the mixture or the mixture may be covered with a heating box.

(Cold rolling process)
In the cold rolling step, the hot-rolled steel sheet after the retention step is recoiled, pickled and cold rolled to obtain a cold-rolled steel sheet.
By performing pickling, oxide scale on the surface of the hot-rolled steel sheet can be removed, and the chemical conversion property and plating property of the cold-rolled steel sheet can be improved. Pickling may be performed under known conditions, and may be performed once or multiple times. The reduction ratio (rolling ratio) of the cold rolling is not particularly limited. For example, it is 20 to 80%.

(Annealing process)
In the annealing step, the cold-rolled steel sheet after the cold rolling step is annealed by holding it at an annealing temperature of 800° C. or more and less than 900° C.
In this annealing process, in the process of heating to the annealing temperature in the austenite single phase region, these precipitates play a role of pinning the prior austenite grain boundaries in the relatively low temperature region. As a result, the precipitates are unevenly distributed on the prior austenite grain boundaries. When further heating is performed and a relatively high temperature region is reached, the precipitates become thermally unstable and dissolve. As a result, the alloying elements can be segregated on the prior austenite grain boundaries.
In this state, by performing rapid cooling as in the post-annealing cooling process described later, the austenite is transformed into martensite, and a high-strength steel plate in which alloy elements are unevenly distributed on the prior austenite grain boundaries can be obtained.
If the annealing temperature is less than 800° C., the amount of austenite produced is small and the dissolution of carbides is insufficient. Therefore, the annealing temperature is set to 800° C. or higher, and preferably 830° C. or higher.
On the other hand, if the annealing temperature is 900° C. or higher, grain growth occurs and the prior austenite grain size becomes coarse, which may suppress segregation of predetermined alloy elements to grain boundaries and deteriorate hydrogen embrittlement resistance.
The holding time at the annealing temperature does not need to be limited, but if the holding time is less than 10 seconds, the austenite fraction 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, but when continuous annealing is assumed, the line length of the equipment will be long, so that the practical upper limit may be about 600 seconds.
The average heating rate up to the annealing temperature is preferably 2 to 35° C./sec.

(Cooling process after annealing)
(Tempering process)
In the post-annealing cooling step, after the annealing step, 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./sec.
This cooling process rapidly cools the steel sheet in a state where alloy elements are segregated at the austenite grain boundaries, and the austenite transforms into martensite, resulting in a steel sheet having a structure mainly composed of martensite and alloy elements segregated at the prior austenite grain boundaries.
If the average cooling rate is less than 20° C./sec, a sufficient amount of martensite is not generated. On the other hand, if the average cooling rate exceeds 100° C./sec, the capacity of the equipment may be insufficient and the equipment may need to be expanded when continuous annealing is assumed, so 100° C./sec is set as the practical upper limit.
Furthermore, if the cooling stop temperature exceeds 300° C., untransformed austenite that has not been transformed into martensite is likely to transform into bainite, which may result in a decrease in strength. On the other hand, even if the cooling stop temperature is less than 25° C., the effect is saturated and a special refrigerant or the like is required, which may result in a decrease in manufacturability and an increase in costs.
The cold-rolled steel sheet after the annealing and cooling process may be further subjected to a tempering process in which the steel sheet is heated to 50 to 550°C and held for 10 to 1000 seconds. By performing this tempering process, alloy elements that are not segregated to the grain boundaries and are present within the grains can be segregated onto the prior austenite grain boundaries. In addition, by turning 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 effect cannot be obtained. On the other hand, if the tempering temperature exceeds 550°C, the strength may decrease due to a decrease in dislocation density in the tempered martensite, leading to a decrease in tensile strength. In addition, carbides may be coarsely precipitated on the prior austenite grain boundaries, deteriorating the hydrogen embrittlement resistance. In addition, if the holding time exceeds 1000 seconds, the strength decreases and the productivity decreases. Tempering may be performed in a continuous annealing facility, or may be performed offline after continuous annealing in a separate facility.

In the above-mentioned post-annealing cooling step, during cooling, holding may be performed in a temperature range of 350 to 650° C. (second temperature range: a temperature range considered to be higher than the Ms point) for 10 to 200 seconds. In this case, the cooling rate 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 step, the steel sheet is cooled from the annealing temperature to a second temperature range of 350 to 650°C at an average cooling rate of 20 to 100°C/s, held in the second temperature range for 10 to 200 seconds, and then cooled from the second temperature range to 25 to 300°C at an average cooling rate of 20 to 100°C/s.
By holding at 350 to 650°C, alloy elements that are not segregated to the grain boundaries and are present within the grains can be segregated onto the prior austenite grain boundaries, improving hydrogen embrittlement resistance. However, if the holding temperature is less than 350°C, bainite transformation is likely to occur, and strength may decrease. If the holding time is less than 20 seconds, the effect of segregating elements present within the grains onto the 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, resulting in a decrease in tensile strength. Furthermore, if the holding time exceeds 200 seconds, bainite transformation and ferrite transformation from austenite are likely to occur. The holding temperature is preferably in the range of 370°C or more and 630°C or less, more preferably 390°C or more and 610°C or less. Furthermore, the holding time is preferably in the range of 30 seconds or more and 180 seconds or less, more preferably 50 seconds or more and 160 seconds or less.
The holding during the cooling step after annealing and the tempering step both promote the segregation of alloying elements on the prior austenite grain boundaries, so either one of these steps may be performed, or both may be performed, or neither step may be performed.

The method for producing a steel sheet according to this embodiment may include a coating layer forming step of forming a coating layer on one or both surfaces of the steel sheet.
The coating layer is preferably a coating layer containing zinc, aluminum, magnesium or an alloy thereof. The coating layer is, for example, a plating layer.
The coating method is not limited, but for example, when forming a coating layer mainly composed of zinc by hot dip plating, a cold rolled steel sheet is adjusted (by heating or cooling) so that the steel sheet temperature is (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 under the conditions exemplified.
The reason why this condition is preferable is that if the temperature of the steel sheet during immersion in the coating bath falls below the hot-dip galvanizing bath temperature -40°C, heat removal during immersion in the coating bath is large and some of the molten zinc may solidify, resulting in deterioration of the coating appearance, whereas if the temperature exceeds the hot-dip galvanizing bath temperature +50°C, operational problems associated with an increase in the coating bath temperature may be induced.
When forming a zinc-based plating layer, the plating bath preferably has an effective Al content (the total Al content minus the total Fe content in the plating bath) of 0.050 to 0.250 mass%, contains Mg as necessary, and the balance is Zn and impurities. If the effective Al content in the plating bath is less than 0.050 mass%, Fe may penetrate excessively into the plating layer, resulting in a decrease in plating adhesion. On the other hand, if the effective Al content in the plating bath exceeds 0.250 mass%, Al-based oxides that inhibit the movement of Fe and Zn atoms may be generated at the boundary between the steel sheet and the plating layer, resulting in a decrease in plating adhesion.

The formation of the coating layer may be performed after the above-mentioned post-annealing cooling step, during the post-annealing cooling step, or during the tempering step. That is, it may be performed as part of the holding at 350 to 650°C during the post-annealing cooling step, or as part of the holding at 50 to 550°C during the tempering step.

When a zinc-based plating layer is formed as the coating layer, an alloying treatment may be further performed. In this case, the steel sheet on which the plating layer is formed may be held at 480 to 550° C. for 1 to 30 seconds, for example.
The alloying step may also be carried out during the post-annealing cooling step or during the tempering step described above, i.e. as part of the holding at 350-650°C during the post-annealing cooling step, or as part of the holding at 50-550°C during the tempering step.

In order to improve paintability and weldability, a top layer of plating may be applied to the surface of the coating layer, or various treatments such as chromate treatment, phosphate treatment, lubricity improvement treatment, weldability improvement treatment, etc. may be applied.

The following is an example of the present invention. The following example is an example of the present invention, and the present invention is not limited to the example described below.

Example 1
Steels having the chemical compositions shown in Tables 1-1 to 1-4 were melted and cast into billets.
The steel slab was inserted into a furnace heated to 1220°C, held there for 60 minutes, and then taken out into the atmosphere and hot rolled to obtain a steel plate (hot-rolled steel plate) having a thickness of 2.8 mm. In the hot rolling, a rolling mill having seven stands was used to continuously perform a total of seven finish rolling passes (so that the inter-pass time was constant), of which four rolling passes had a rolling reduction of more than 20%. In addition, the inter-pass time between each rolling pass that gave a rolling reduction of 20% or more in the finish rolling and the rolling pass immediately preceding each of the rolling passes was 0.6 seconds. The start temperature of the finish rolling was 1060°C, and the end temperature was 870°C.
For this hot-rolled steel sheet, water cooling was started 2.2 seconds after the end of hot rolling, and the sheet was cooled to 580°C at an average cooling rate of 38.0°C/s and coiled, and then loaded into a furnace at 530°C and held there for 1,800 seconds.
Subsequently, the oxide scale of this hot-rolled steel sheet was removed by pickling, and the hot-rolled steel sheet was subjected to cold rolling at a rolling reduction of 50.0% to obtain a cold-rolled steel sheet having a thickness of 1.4 mm.
This cold-rolled steel sheet was heated to 880°C at an average heating rate of 12.0°C/s, held at 880°C for 120 seconds, and then cooled to 150°C at an average cooling rate of 42.0°C/s.
Thereafter, the cold-rolled steel sheet was subjected to tempering by reheating to 230° C. and holding for 180 seconds. No plating treatment was performed.

The chemical composition was analyzed using samples taken from the obtained steel plate. The results were the same as the chemical composition of the steel shown in Tables 1-1 to 1-4.

In addition, for the obtained cold-rolled steel sheets, the area ratios of ferrite, martensite, tempered martensite, and the remainder (one or more of bainite, pearlite, and retained austenite) in the microstructure of the t/4 portion were determined by the method described above.
The results are shown in Table 2.

Furthermore, the concentration of each alloy element on the austenite grain boundary was measured by the method described above, and E _GB was calculated.

In addition, the tensile strength, total elongation, and hydrogen embrittlement resistance (hydrogen embrittlement resistance characteristics) of the obtained cold-rolled steel sheets were evaluated using the following methods.

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

(Method of evaluating hydrogen embrittlement resistance)
The hydrogen embrittlement resistance of the hot-dip galvanized steel sheet manufactured by the steel sheet manufacturing method according to the embodiment of the present invention was evaluated by the following method. Specifically, the steel sheet was sheared at a clearance of 15%, and then a U-bending test was performed at 8R. A strain gauge was attached to the center of the obtained test piece, and both ends of the test piece were tightened with bolts to apply stress. The applied stress was calculated from the strain of the monitored strain gauge. The applied load stress was 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 x 0.8 = 1212 MPa). This is because the residual stress introduced during forming is considered to correspond to the tensile strength of the steel sheet.
The obtained U-bend test piece was immersed in an HCl aqueous solution with a pH of 2 at a liquid temperature of 25°C and held for 96 hours to check for the presence or absence of cracks. The lower the pH of the HCl aqueous solution and the longer the immersion time, the greater the amount of hydrogen that penetrates into the steel sheet, resulting in a severe hydrogen embrittlement environment. After immersion, the U-bend test piece was rated as NG if a crack longer than 1.00 mm was observed, and as OK if no crack longer than 1.00 mm was observed. The OK evaluation was deemed to have passed, and the NG evaluation was deemed to have failed.

Steel plates with a tensile strength of 1,500 MPa or more and an evaluation of hydrogen embrittlement resistance that is OK were evaluated as having high strength and excellent hydrogen embrittlement resistance.

Referring to Tables 1-1 and 2, Production No. P had a tensile strength of less than 1500 MPa due to the low C content.
Production No. Q had a high C content and therefore had poor hydrogen embrittlement resistance.
Since Production No. R had a high Si content, the precipitation of alloy carbides was suppressed in the hot rolling process, 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.
Production No. S had a tensile strength of less than 1500 MPa due to the low Mn content.
Production No. T had a high Mn content and thus a low _EGB , and therefore had low hydrogen embrittlement resistance.
Production No. U had a high P content, and therefore the _EGB and hydrogen embrittlement resistance were reduced.
Production No. V had a high S content, and therefore had poor hydrogen embrittlement resistance.
Production No. W had a high N content, which resulted in the formation of coarse nitrides, resulting in a decrease in hydrogen embrittlement resistance.
Since Production No. X had a high Al content, coarse Al oxides were formed, resulting in a decrease in hydrogen embrittlement resistance.
Production No. Y had a high B content, which resulted in the formation of coarse B oxides, resulting in a decrease in hydrogen embrittlement resistance.
Production No. Z had a high Ti content, which resulted in the formation of coarse carbonitrides, which reduced hydrogen embrittlement resistance. In addition, the formation of coarse carbonitrides reduced the amount of Ti segregated at grain boundaries, which reduced _EGB . In addition, the amount of C, which is effective in improving strength, was reduced, resulting in a tensile strength of 1500 MPa or less.
Since Production No. AA had a high Nb content, coarse Nb carbides were formed, resulting in a decrease in hydrogen embrittlement resistance.
Since Production No. AB had a high V content, coarse V carbides were formed, and the hydrogen embrittlement resistance was reduced.
Since Production No. AC had a high O content, oxides were formed, resulting in a decrease in hydrogen embrittlement resistance.
In Production No. AD, the Mo content was high, which resulted in a large amount of carbonitride precipitation and a decrease in hydrogen embrittlement resistance. In addition, the formation of coarse carbonitrides reduced the amount of Mo segregated at the grain boundaries, resulting in a decrease in _EGB .
In Production No. AE, the Cr content was high, so coarse Cr carbides were formed at the center segregation position in the steel, which reduced the hydrogen embrittlement resistance. In addition, the formation of coarse carbonitrides reduced the amount of Cr segregated at the grain boundaries, which reduced _EGB .
In Production No. AF, the Co content was high, so coarse Co carbides were formed, which reduced hydrogen embrittlement resistance. In addition, the formation of coarse carbonitrides reduced the amount of Co segregated at the grain boundaries, which reduced _EGB .
Production No. AG had a high Ni content, and therefore had poor hydrogen embrittlement resistance.
Since Production No. AH had a high Cu content, the EGB _value fell to less than 0.50, resulting in a decline in hydrogen embrittlement resistance.
In Production No. AI, the W content was high, so coarse W precipitates were formed, and the hydrogen embrittlement resistance was reduced. In addition, the amount of W effective for strengthening grain boundaries was reduced, and _EGB was less than 0.50, which reduced the hydrogen embrittlement resistance.
Since Production No. AJ had a high Ta content, a large number of fine Ta carbides were precipitated, and the hydrogen embrittlement resistance was reduced.
Production No. AK had a high Sn content, which caused grain boundary embrittlement, resulting in an EGB _value below 0.50, and therefore a decrease in hydrogen embrittlement resistance.
Production Nos. AL and AM had high Sb and As contents, respectively, so that _EGB fell below 0.50 due to grain boundary segregation, and hydrogen embrittlement resistance was reduced.
Production Nos. AN and AO had high Mg and Ca contents, respectively, which resulted in the formation of coarse inclusions, which reduced the hydrogen embrittlement resistance.
Since Production Nos. AP to AS had high contents of Y, Zr, La and Ce, respectively, coarse oxides were formed and the hydrogen embrittlement resistance was reduced.
In contrast, in Production Nos. A to O, by appropriately controlling the chemical composition and structure of the steel plate and the grain boundary strength _EGB of the prior austenite grain boundary, it was possible to obtain steel plates having high strength and excellent resistance to hydrogen embrittlement.

Example 2
Furthermore, in order to investigate the influence of the manufacturing conditions, the steel types (steel Nos. A to O) in which excellent properties were recognized in Example 1 were targeted, and in the same equipment as in Example 1, the steel pieces were inserted into a furnace heated to 1250 to 1100°C, held for 60 minutes, and then taken out into the atmosphere, and hot-rolled steel sheets with a sheet 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 under the conditions described in Tables 3-1 to 3-4 after coiling. A part of the cold-rolled steel sheets was used as a plated steel sheet on which a plating layer was formed. Here, the symbols GI and GA of the plating treatment indicate the method of the zinc plating treatment, and GI is a steel sheet in which a zinc plating layer was formed on the surface of the steel sheet by immersing the steel sheet in a hot-dip galvanizing bath at 460°C, and GA is a steel sheet in which an alloy layer of iron and zinc was formed on the surface of the steel sheet by immersing the steel sheet in a hot-dip galvanizing bath and then raising the temperature to 485°C. When intermediate holding was performed in the second temperature range, plating was performed as it was after the intermediate holding (without cooling to room temperature). When intermediate holding was not performed in the second temperature range, plating was performed during cooling to 25 to 300°C. In addition, in Tables 3-3 to 3-4, examples in which tempering is described as "-" are examples in which tempering was not performed. In addition, the interpass times in the tables are the interpass times for each pass with a reduction rate of 20% or more (because rolling was performed with a tandem rolling mill, each interpass time was equal). In addition, in Tables 3-3 and 3-4, the holding time in the cooling step after annealing is the holding time in the second temperature range when cooling was performed to the second temperature range, but when the cooling stop temperature is outside the second temperature range, it is the holding time around that temperature.

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

In addition, the tensile strength and total elongation of the obtained cold-rolled steel sheet were evaluated in the same manner as in Example 1.

The hydrogen embrittlement resistance was evaluated by the following method.
(Method of evaluating hydrogen embrittlement resistance)
The hydrogen embrittlement resistance of the hot-dip galvanized steel sheet manufactured by the steel sheet manufacturing method according to the embodiment of the present invention was evaluated by the following method. Specifically, the steel sheet was sheared at a clearance of 15%, and then a U-bending test was performed at 8R. A strain gauge was attached to the center of the obtained test piece, and both ends of the test piece were tightened with bolts to apply stress. The applied stress was calculated from the strain of the monitored strain gauge. The applied load stress was a stress corresponding to 80% of the tensile strength (TS) (for example, in the case of A-1 in Table 4, the applied stress = 1540 MPa x 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-bend test piece was immersed in an HCl aqueous solution with a pH of 2 at a liquid temperature of 25°C and held for 96 hours to check for the presence or absence of cracks. The lower the pH of the HCl aqueous solution and the longer the immersion time, the greater the amount of hydrogen that penetrates into the steel sheet, resulting in a severe hydrogen embrittlement environment. After immersion, the total length of the cracks in the U-bend test piece was measured (if multiple cracks were found, the total length was taken as the sum of the values measured individually).
The smaller the total length of the cracks, the better the hydrogen embrittlement resistance, but in particular, cracks longer than 1.00 mm were rated as NG, no cracks or minor cracks with a length of 1.00 mm or less were rated as OK, and among the OK grades, no cracks or cracks with a length of 0.70 mm or less were rated as Ex. Grades of OK and Ex were considered to be pass, and grades of NG were considered to be fail.
The results are shown in Table 4.

Referring to Table 4, in 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 the precipitation of alloy carbides was delayed. As a result, the alloy elements that contribute to improving the grain boundary strength could not be segregated to the grain boundaries, and the _EGB was below 0.50. As a result, the hydrogen embrittlement resistance was reduced.
In Production No. B-2, the rolling start temperature of the hot rolling was low, so that unrecrystallized austenite remained, the precipitation of alloy carbides was delayed, and the _EGB was below 0.50. As a result, the hydrogen embrittlement resistance was reduced.
In Production No. C-2, the rolling start temperature of the hot rolling was high, so that the crystal grains of the recrystallized austenite became coarse, the precipitation of alloy carbides was delayed, and the _EGB was below 0.50. As a result, the hydrogen embrittlement resistance was reduced.
In Production No. D-2, the rolling end temperature was low, so that unrecrystallized austenite remained, the precipitation of alloy carbides was delayed, and _EGB was below 0.50. As a result, the hydrogen embrittlement resistance was reduced.
In Production No. E-2, the rolling end temperature was high, so that the crystal grains of the recrystallized austenite became coarse, the precipitation of alloy carbides was delayed, and the _EGB fell below 0.50. As a result, the hydrogen embrittlement resistance was reduced.
In Production No. F-2, the coiling temperature was low, which delayed the diffusion of alloy elements and inhibited the precipitation of alloy carbides, so that _EGB was below 0.50. As a result, hydrogen embrittlement resistance was reduced.
In Production No. G-2, the coiling temperature was high, so that an internal oxide layer was formed on the surface layer of the hot-rolled steel sheet, and cracks occurred on the surface of the steel sheet during subsequent processing. Therefore, analysis of the structure and evaluation of the mechanical properties were not performed.
In Production No. H-2, the average cooling rate to the coiling temperature was low, so that ferrite and pearlite transformation occurred, the precipitation of alloy carbides was suppressed, and the _EGB was below 0.50. As a result, the hydrogen embrittlement resistance was reduced.
In Production No. K-2, the interpass time in the finish rolling with a reduction rate of 20% or more was long, which delayed the precipitation of alloy carbides in the hot rolling process. As a result, _EGB was below 0.50, and hydrogen embrittlement resistance was reduced.
In Production No. L-2, the holding temperature after the hot rolling process was low, and the retention time at 400 to 550°C was less than 600 seconds, so that the precipitation of alloy carbides did not occur sufficiently, and the _EGB was less than 0.50. As a result, the hydrogen embrittlement resistance was deteriorated.
In Production No. M-2, the holding temperature after the hot rolling process was high, and the retention time at 400 to 550°C was less than 600 seconds, so that the precipitation of alloy carbides did not occur sufficiently, and the _EGB was below 0.50. As a result, the hydrogen embrittlement resistance was reduced.

In Production No. A-3, the annealing temperature was low, so that ferrite transformation progressed during holding, and the tensile strength fell below 1500 MPa.
In Production No. B-3, the annealing temperature was high, so the concentration of alloy elements segregating at the grain boundaries decreased. As a result, _EGB was below 0.50, and hydrogen embrittlement resistance was reduced.
In Production No. C-3, the average cooling rate from the annealing temperature was small, so ferrite transformation occurred during the cooling process and the tensile strength did not reach 1,500 MPa.
In Production No. F-3, the cooling stop temperature was high, so bainite transformation occurred and the tensile strength did not reach 1500 MPa.
In Production No. H-3, the tempering temperature was high, so that softening of martensite occurred and the tensile strength did not reach 1500 MPa.
In Production No. J-3, the tempering time was long, so that the martensite was softened excessively and the tensile strength did not reach 1500 MPa.
In Production No. K-3, the cooling stop temperature and intermediate holding temperature in the cooling step after annealing were low (outside the second temperature range), so bainite transformation occurred and the tensile strength did not reach 1500 MPa.
In Production No. L-3, the cooling stop temperature and intermediate holding temperature in the cooling step after annealing were high (outside the second temperature range), so ferrite transformation and pearlite transformation occurred, and the tensile strength did not reach 1500 MPa.
In Production No. M-3, the holding time at the second cooling stop temperature was long, so bainite transformation occurred during holding, and the tensile strength did not reach 1,500 MPa.
In Production No. N-3, the cooling rate from the second cooling stop temperature was small, so ferrite transformation and bainite transformation occurred during the cooling process, and the tensile strength did not reach 1,500 MPa.

In contrast, in all examples of the present invention, by appropriately controlling hot rolling, coiling, annealing, etc., steel sheets with high strength and excellent resistance to hydrogen embrittlement were obtained.

Figure 1 is a diagram showing the relationship between _EGB and tensile strength on the hydrogen embrittlement resistance of steel sheets in Examples 1 and 2. In Figure 1, ▲ indicates an example in which the hydrogen embrittlement resistance target was not achieved, and ◯ indicates an example in which the hydrogen embrittlement resistance target was achieved. As shown in Figure 1, by making _EGB 0.50 or more, it is possible to obtain excellent hydrogen embrittlement resistance even in high-strength materials of 1500 MPa or more.

According to the present invention, it is possible to provide a high-strength steel sheet with excellent resistance to hydrogen embrittlement. When this steel sheet is applied to automobile steel sheets, it contributes to reducing the weight of the vehicle body and improving fuel efficiency.

Claims (3)

In mass percent,
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 to 0.0200%,
O: 0.0001-0.0200%,
Co: 0 to 0.500%,
Ni: 0-1.000%,
Mo: 0-1.000%,
Cr: 0-2.000%,
Ti: 0 to 0.500%,
B: 0 to 0.0100%,
Nb: 0 to 0.500%,
V: 0 to 0.500%,
Cu: 0-0.500%,
W: 0-0.100%,
Ta: 0-0.100%,
Mg: 0 to 0.050%,
Ca: 0-0.050%,
Y: 0 to 0.050%,
Zr: 0 to 0.050%,
La: 0 to 0.050%,
Ce: 0 to 0.050%,
Sn: 0 to 0.050%,
Sb: 0 to 0.050%,
A chemical composition consisting of As: 0 to 0.050%, and the balance: Fe and impurities,
The microstructure, in terms of 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;
When an interface between adjacent martensite and tempered martensite with an orientation difference of 15 degrees or more is defined as a prior austenite grain boundary, the bond strength energy E _GB determined by the concentration of each alloy element on the prior austenite grain boundary satisfies the following formula (1):
The tensile strength is 1500 MPa or more.
A steel plate characterized by:
_EGB =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, the [chemical symbols] in the formula represent the concentrations of the respective alloying elements in mass % on the prior austenite grain boundaries. The chemical composition is
Co: 0.01 to 0.500%,
Ni: 0.01-1.000%,
Mo: 0.01-1.000%,
Cr: 0.001-2.000%,
Ti: 0.001 to 0.500%,
B: 0.0001 to 0.0100%,
Nb: 0.001-0.500%,
V: 0.001-0.500%,
Cu: 0.001 to 0.500%,
W: 0.001-0.100%,
Ta: 0.001 to 0.100%,
Mg: 0.001-0.050%,
Ca: 0.001-0.050%,
Y: 0.001-0.050%,
Zr: 0.001 to 0.050%,
La: 0.001 to 0.050%,
Ce: 0.001 to 0.050%,
Sn: 0.001 to 0.050%,
Sb: 0.001 to 0.050%, and As: 0.001 to 0.050%,
Contains one or more selected from the group consisting of
2. The steel sheet according to claim 1 , A coating layer containing zinc, aluminum, magnesium or an alloy thereof is provided on the surface.
The steel sheet according to claim 1 or 2, characterized in that

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