CN118140000A

CN118140000A - Steel plate

Info

Publication number: CN118140000A
Application number: CN202280070217.2A
Authority: CN
Inventors: 大贺光阳; 竹田健悟; 中野克哉
Original assignee: Nippon Steel and Sumitomo Metal Corp
Current assignee: Nippon Steel Corp
Priority date: 2021-10-21
Filing date: 2022-10-21
Publication date: 2024-06-04
Also published as: WO2023068369A1; KR20240069758A; JPWO2023068369A1

Abstract

The present invention provides a steel sheet having a predetermined chemical composition, wherein the microstructure comprises ferrite in terms of area ratio: less than 5.0%, martensite and tempered martensite: aggregate over 90.0%, remainder: when the boundary between adjacent martensite and tempered martensite having a difference in orientation of 15 ° or more is a prior austenite grain boundary, 1 or 2 or more of bainite, pearlite, and retained austenite has a bonding strength energy E _GB determined by the concentration of each alloy element in the prior austenite grain boundary of 0.50 or more and a tensile strength of 1500MPa or more.

Description

Steel plate

Technical Field

The present invention relates to a steel sheet.

The present application claims priority based on japanese patent application publication No. 2021-172424, filed on 21 of 10 in 2021, the contents of which are incorporated herein by reference.

Background

Today, which is highly specialized in industrial technology, special and high performance is required for materials used in each technical field. In particular, in regard to steel sheets for automobiles, there is a significant increase in demand for high-strength steel sheets in order to improve fuel efficiency by reducing the weight of the automobile body in view of global environment. However, most metal materials are associated with a high strength, and these materials deteriorate various characteristics, particularly an increase in hydrogen embrittlement sensitivity. In steel members, if the tensile strength is 1200MPa or more, particularly, the hydrogen embrittlement sensitivity is improved, and it is known that in bolt steels which have been advanced to have higher strength in the automotive field, there are cases of hydrogen embrittlement cracking. Therefore, a high-strength steel sheet having a tensile strength of 1500MPa or more is strongly demanded to fundamentally solve the problem of hydrogen embrittlement.

In a high-strength steel sheet having a tensile strength of 1500MPa or more, the microstructure mainly includes martensite and tempered martensite, but in such a high-strength steel sheet, hydrogen that has entered the steel is segregated into the grain boundaries of martensite, and the grain boundaries are embrittled (the grain boundary strength is lowered), so that cracks (hydrogen embrittlement) occur. Since hydrogen intrusion occurs even at room temperature, there is no method of perfectly suppressing hydrogen intrusion, and in order to fundamentally solve this problem, improvement of the steel internal structure is necessary.

Conventionally, various techniques for improving the hydrogen embrittlement resistance (also referred to as hydrogen embrittlement resistance) of a high-strength steel sheet have been proposed (for example, see patent documents 1 to 6).

Patent document 1 discloses an ultra-high strength steel sheet excellent in hydrogen embrittlement resistance, which is characterized in that: as an ultra-high strength steel sheet excellent in hydrogen embrittlement resistance and workability, C: more than 0.25 to 0.60 percent of Si:1.0 to 3.0 percent of Mn:1.0 to 3.5 percent of P:0.15% or less, S: less than 0.02%, al: less than 1.5% (excluding 0%), mo: less than 1.0% (excluding 0%), nb:0.1% or less (excluding 0%), the remainder including iron and unavoidable impurities, and the metal structure after the drawing work having a working rate of 3% satisfying the retained austenite structure in terms of area ratio relative to the whole structure: more than 1 percent of bainitic ferrite and martensite: total 80% or more ferrite and pearlite: the total of the retained austenite grains is 9% or less (including 0%), and the average axial ratio (major axis/minor axis) of the retained austenite grains is satisfied: 5 or more, and the tensile strength is 1180MPa or more.

Patent document 2 discloses a high-strength steel sheet having a tensile strength of 1500MPa or more, which is obtained by including si+mn in the steel component: 1.0% or more, forming a layer of ferrite and carbide in a main phase structure, and further forming a layer structure of carbide having an aspect ratio of 10 or more and a spacing of 50nm or less between the layers to a volume ratio of 65% or more relative to the entire structure, and defining a fraction of carbide having an aspect ratio of 10 or more and an angle of 25 ° or less relative to a rolling direction to 75% or more in terms of an area ratio, thereby making the bending property in the rolling direction and the delayed fracture resistance excellent.

Patent document 3 discloses an ultra-high strength cold-rolled steel sheet excellent in bendability, which is characterized in that: a thin ultra-high strength cold-rolled steel sheet excellent in bending properties and delayed fracture resistance comprises, in mass%, C:0.15 to 0.30 percent of Si:0.01 to 1.8 percent of Mn:1.5 to 3.0 percent of P: less than 0.05%, S: less than 0.005%, al: 0.005-0.05%, N:0.005% or less, the remainder comprising Fe and unavoidable impurities, the steel sheet having a steel sheet surface soft portion satisfying the relationship of "hardness of the steel sheet surface soft portion/hardness of the steel sheet central portion is 0.8 or less", the steel sheet surface soft portion having a ratio of 0.10 or more to 0.30 or less in sheet thickness, and tempered martensite in the steel sheet surface soft portion being 90% or more by volume, the steel sheet central portion having a structure of tempered martensite, and tensile strength of 1270MPa or more.

Patent document 4 discloses a cold-rolled steel sheet having a tensile strength of 1470MPa or more and excellent bending workability and delayed fracture resistance, which contains, in mass%, C:0.15 to 0.20 percent of Si:1.0 to 2.0 percent of Mn:1.5 to 2.5 percent of P: less than 0.020%, S: less than 0.005%, al:0.01 to 0.05 percent of N: less than 0.005%, ti: less than 0.1%, nb:0.1% or less, B:5 to 30ppm, the remainder comprising Fe and unavoidable impurities, and the tempered martensite phase being 97% or more by volume and the retained austenite phase being less than 3% by volume.

Patent document 5 discloses an ultra-high strength steel sheet having a tensile strength of 1470MPa or more, which exhibits excellent delayed fracture resistance even at the cut end, and which has a composition containing, in mass%, C:0.15 to 0.4 percent of Mn:0.5 to 3.0 percent of Al: 0.001-0.10%, the remainder comprising iron and unavoidable impurities, wherein P, S, N is limited to P:0.1% or less, S: less than 0.01%, N:0.01% or less, and a structure including martensite in an area ratio relative to the entire structure: 90% or more of retained austenite: 0.5% or more, and the area ratio of the region where the local Mn concentration is 1.1 times or more of the Mn content of the whole steel sheet is 2% or more, and the tensile strength is 1470MPa or more.

Patent document 6 discloses an ultra-high strength cold-rolled steel sheet having excellent hydrogen embrittlement resistance and a tensile strength of 1300MPa or more, which comprises C:0.150 to 0.300 percent of Si:0.001 to 2.0 percent of Mn:2.10 to 4.0 percent of P: less than 0.05%, S: less than 0.01%, N: less than 0.01%, al:0.001% -1.0%, ti:0.001% -0.10%, B: the values of the solid solution B content solB [ mass% ] and the prior austenite grain diameter Dgamma [ mu ] m satisfy the relation of solB.Dgamma.gtoreq.0.0010, wherein the polygonal ferrite in the steel structure is 10% or less, the bainite is 30% or less, the retained austenite is 6% or less, the tempered martensite is 60% or more, the number density of Fe carbide in the tempered martensite is 1 x 10 ⁶/mm² or more, the average dislocation density of the whole steel is 1.0 x 10 ¹⁵～2.0×10¹⁶/m², and the crystal grain diameter is 7.0 [ mu ] m or less.

Prior art literature

Patent literature

Patent document 1: japanese patent laid-open No. 2006-207019

Patent document 2: japanese patent application laid-open No. 2010-138489

Patent document 3: japanese patent laid-open publication No. 2011-179030

Patent document 4: japanese patent application laid-open No. 2010-215958

Patent document 5: japanese patent laid-open publication 2016-153524

Patent document 6: japanese patent laid-open publication 2016-050343

Disclosure of Invention

Problems to be solved by the invention

As described above, several techniques have been proposed for improving the hydrogen embrittlement resistance (hydrogen embrittlement resistance) of high-strength steel sheets. However, patent document 1 discloses only hydrogen embrittlement resistance when a stress of 1000MPa is applied, and therefore does not show any technical solution for hydrogen embrittlement resistance when a higher stress is applied.

As described above, hydrogen embrittlement occurs because hydrogen is accumulated in grain boundaries and the bonding strength of the grain boundaries is reduced. Therefore, it is considered that if the bonding strength of the grain boundaries can be improved, cracking due to hydrogen embrittlement can be suppressed. However, patent documents 1 to 6 do not disclose a method for improving the hydrogen embrittlement resistance from such a viewpoint. In recent years, the requirements for hydrogen embrittlement resistance have become more stringent, and in patent documents 1 to 6, such stringent requirements may not be satisfied.

That is, conventionally, there is room for improvement in hydrogen embrittlement resistance of a high-strength steel sheet having a microstructure mainly composed of martensite and tempered martensite.

In patent document 2, the steel sheet has a structure in which a pearlite structure is a main phase, a ferrite phase in a remaining portion of the structure is 20% or less by volume ratio with respect to the entire structure, and a platelet spacing of the pearlite structure is 500nm or less, and the rolling ratio is set to a steel sheet having a vickers hardness of HV200 or more: and 60% or more (preferably 75% or more) of the cold rolling. Therefore, it can be easily inferred that the anisotropy is strong and the cold formability of the member is low.

In addition, in patent document 3, in order to improve the delayed fracture characteristics, it is necessary to maintain the temperature at 650 ℃ or 700 ℃ for 20 minutes or more in an atmosphere having a dew point of 15 ℃ or more, and there is also a problem of low productivity.

The present invention has been made in view of the above problems. The subject of the invention is that: provided is a steel sheet having excellent hydrogen embrittlement resistance on the premise of having a microstructure mainly composed of martensite and tempered martensite.

Means for solving the problems

As described above, hydrogen embrittlement is considered to be cracking that occurs with the grain boundaries as the starting point, because hydrogen in steel segregates to the grain boundaries, reducing the bonding strength of the grain boundaries. Accordingly, the present inventors have made various studies on a method for improving hydrogen embrittlement resistance, focusing on the bonding strength of grain boundaries.

As a result, the present inventors found that: by segregating the predetermined alloy element to the grain boundaries, the bonding strength of the grain boundaries can be improved, and the intruded hydrogen is made difficult to segregate to the grain boundaries, so that even if the hydrogen intrudes, the decline of the bonding strength of the grain boundaries due to the hydrogen can be suppressed.

The present invention has been made in view of the above-described findings. The gist of the present invention is as follows.

[1] One embodiment of the present invention relates to a steel sheet having a chemical composition comprising C：0.150～0.400％、Si：0.01～2.00％、Mn：0.80～2.00％、P：0.0001～0.0200％、S：0.0001～0.0200％、Al：0.001～1.000％、N：0.0001～0.0200％、O：0.0001～0.0200％、Co：0～0.500％、Ni：0～1.000％、Mo：0～1.000％、Cr：0～2.000％、Ti：0～0.500％、B：0～0.0100％、Nb：0～0.500％、V：0～0.500％、Cu：0～0.500％、W：0～0.100％、Ta：0～0.100％、Mg：0～0.050％、Ca：0～0.050％、Y：0～0.050％、Zr：0～0.050％、La：0～0.050％、Ce：0～0.050％、Sn：0～0.050％、Sb：0～0.050％、As：0～0.050％ in mass% and the remainder: fe and impurities, the microstructure comprising ferrite in terms of area ratio: less than 5.0%, martensite and tempered martensite: aggregate over 90.0%, remainder: when an interface having a difference (orientation difference) between orientation of adjacent martensite and tempered martensite of 15 ° or more is used as a prior austenite grain boundary, 1 or more of bainite, pearlite, and retained austenite, the bonding strength energy E _GB determined by the concentration of each alloy element in the prior austenite grain boundary satisfies the following formula (1), and the tensile strength is 1500MPa or more;

E_GB＝1+(3×[Co]+0.7×[Ni]+5.5×[Mo]+0.7×[Cr]+2.9×[Ti]+47×[B]+4.3×[Nb]+4.5×[V]+5.2×[W]+3.1×[Ta]+4.3×[Zr]-0.25×[Mn]-0.1×[P]-[Cu]-1.1×[Sn]-0.6×[Sb]-0.9×[As])≥0.50 (1)

Wherein [ chemical symbol ] in the formula represents the concentration of each alloy element in mass% on the prior austenite grain boundary.

[2] The steel sheet according to the above [1], wherein the chemical composition may also contain a composition selected from Co：0.01～0.500％、Ni：0.01～1.000％、Mo：0.01～1.000％、Cr：0.001～2.000％、Ti：0.001～0.500％、B：0.0001～0.0100％、Nb：0.001～0.500％、V：0.001～0.500％、Cu：0.001～0.500％、W：0.001～0.100％、Ta：0.001～0.100％、Mg：0.001～0.050％、Ca：0.001～0.050％、Y：0.001～0.050％、Zr：0.001～0.050％、La：0.001～0.050％、Ce：0.001～0.050％、Sn：0.001～0.050％、Sb：0.001～0.050％ and As:0.001 to 0.050% of 1 or more than 2 kinds.

[3] The steel sheet according to the above [1] or [2], wherein a coating layer containing zinc, aluminum, magnesium or an alloy thereof may be provided on the surface.

Effects of the invention

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

Drawings

FIG. 1 is a graph showing the relationship between the hydrogen embrittlement resistance of a steel sheet and E _GB and tensile strength in examples of the present invention.

Detailed Description

A steel sheet according to an embodiment of the present invention (steel sheet according to the present embodiment) will be described below.

The steel sheet according to the present embodiment has a predetermined chemical composition,

The microstructure comprises ferrite in terms of area ratio: less than 5.0%, martensite and tempered martensite: aggregate over 90.0%, remainder: 1 or more than 2 kinds of bainite, pearlite and retained austenite,

When the interface between adjacent martensite and tempered martensite having a difference in orientation of 15 DEG or more is defined as a prior austenite grain boundary (prior γ grain boundary), the bonding strength energy E _GB determined by the concentration of each alloy element in the prior austenite grain boundary is 0.50 or more,

The tensile strength is more than 1500 MPa.

< Chemical composition >

First, the ranges of the contents of the elements constituting the chemical composition of the steel sheet according to the present embodiment will be described. Hereinafter, "%" related to the content of an element means "% by mass". The range indicated by the "to" includes the values at both ends as the lower limit or the upper limit.

C：0.150～0.400％

C is an element effective for inexpensively increasing the tensile strength. If the C content is less than 0.150%, the targeted tensile strength is not obtained, and the fatigue properties of the weld zone are deteriorated. Therefore, the C content is set to 0.150% or more. The C content may be 0.160% or more, 0.180% or more, or 0.200% or more.

On the other hand, 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％

Si functions as a deoxidizer and is an element that affects the morphology of carbides and residual austenite after heat treatment. If the Si content is less than 0.01%, it becomes difficult to suppress the formation of coarse oxides. The coarse oxide becomes a starting point of cracks, and the cracks propagate in the steel material, thereby deteriorating hydrogen embrittlement resistance. Therefore, the Si content is set to 0.01% or more. The Si content may be 0.05% or more, 0.10% or more, or 0.30% or more.

On the other hand, if the Si content exceeds 2.00%, precipitation of alloy carbide 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 for improving the strength of the steel sheet. If the Mn content is less than 0.80%, the effect cannot be obtained sufficiently. Therefore, the Mn content is set to 0.80% or more. The Mn content may be 1.00% or more, or 1.20% or more.

On the other hand, if the Mn content exceeds 2.00%, mn not only promotes co-segregation with P, S, but also sometimes deteriorates 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 to ferrite grain boundaries and promotes grain boundary embrittlement. If the P content exceeds 0.0200%, the hydrogen embrittlement resistance is significantly reduced by grain boundary embrittlement. Therefore, the P content is set to 0.0200% or less. The P content may be 0.0180% or less, 0.0150% or less, or 0.0120% or less.

The smaller the P content, the more preferable. However, when the P content is set to less than 0.0001%, the time required for refining increases, thereby incurring a significant increase in cost. Therefore, the P content is set to 0.0001% or more. The P content may be 0.0005% or more, 0.0010% or more, or 0.0020% or more.

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 crack initiation points during cold working becomes remarkable. In this case, even if the grain boundaries are strengthened, cracks from nonmetallic inclusions occur, and the cracks propagate in the steel material, thereby deteriorating hydrogen embrittlement resistance. Therefore, the S content is set to 0.0200% or less. The S content may be 0.0180% or less, 0.0150% or less, or 0.0120% or less.

The smaller the S content, the more preferable. However, when the S content is set to less than 0.0001%, the time required for refining increases, which results in a significant increase in cost. Therefore, the S content is set to 0.0001% or more. The S content may be 0.0005% or more, 0.0010% or more, or 0.0020% or more.

Al：0.001～1.000％

Al acts as a deoxidizer for steel, and stabilizes ferrite. If the Al content is less than 0.001%, a sufficient effect cannot be obtained. Therefore, the Al content is set to 0.001% or more. The Al content may be 0.005% or more, 0.010% or more, or 0.020% or more.

On the other hand, if the Al content exceeds 1.000%, coarse Al oxide is formed. The coarse oxide becomes a crack initiation point. Therefore, if coarse Al oxide is formed, cracks are generated by the coarse Al oxide even if the grain boundaries are reinforced, and the cracks propagate in the steel material, thereby deteriorating the hydrogen embrittlement resistance. Therefore, the Al content is set to 1.000% or less. The Al content may be 0.950% or less, 0.900% or less, or 0.800% or less.

N：0.0001～0.0200％

N is an element that forms coarse nitrides in the steel sheet and reduces the hydrogen embrittlement resistance of the steel sheet. N is an element that causes blowholes during welding.

If the N content exceeds 0.0200%, the occurrence of pores becomes remarkable while deteriorating the hydrogen embrittlement resistance. Therefore, the N content is set to 0.0200% or less. The N content may be 0.0180% or less, 0.0160% or less, or 0.0120% or less.

On the other hand, when the N content is set to less than 0.0001%, the manufacturing cost increases greatly. 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 an oxide and deteriorates hydrogen embrittlement resistance. In particular, since oxides are often present in the form of inclusions, if they are present in the punched end face or the cut face, scratches in the form of notches and coarse dimples (small) are formed in the end face, stress concentration is caused during the working, and the oxides become the origin of cracks, which greatly deteriorate the workability. If the O content exceeds 0.0200%, the above-mentioned workability tends to be remarkable. Therefore, the O content is set to 0.0200% or less. The O content may be 0.0180% or less, 0.0150% or less, or 0.0100% or less.

Preferably, the O content is relatively low. However, setting the O content to less than 0.0001% is economically undesirable because it will incur excessive increase in cost. Therefore, the O content is set to 0.0001% or more. The O content may be 0.0005% or more, 0.0010% or more, or 0.0015% or more.

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 contains the above elements, and the remainder may include 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, as as an optional component in place of part of Fe in the remaining part in order to improve various properties.

These elements may not be necessarily contained, and therefore the lower limit thereof is 0%. Further, even if the following elements are contained as impurities, the effects of the steel sheet according to the present embodiment are not impaired.

Co：0～0.500％

Co is an element effective for controlling the morphology of carbide and improving the strength of steel sheet. In addition, co is an element that contributes to improvement of grain boundary bonding strength. Therefore, co may be contained. In order to obtain a sufficient effect, the Co content is preferably set to 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 carbide precipitates. In this case, hydrogen embrittlement resistance may be deteriorated. Therefore, the Co content is set to 0.500% or less. The Co content may be 0.450% or less, 0.400% or less, or 0.300% or less.

Ni：0～1.000％

Ni is an element effective for improving the strength of the steel sheet. In addition, ni is also an element that contributes to improving the grain boundary bonding strength. Ni is also an element effective in improving wettability and promoting alloying reaction. Therefore, ni may be contained. In order to obtain the above-described effects, the Ni content is preferably set to 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 be lowered. 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 for improving the strength of the steel sheet. Mo is an element effective for suppressing ferrite transformation generated when heat treatment is performed by a continuous annealing apparatus or a continuous hot dip galvanization apparatus. In addition, mo is an element that contributes to improvement of grain boundary bonding strength. Therefore, mo may be contained. In order to obtain the above-described effects, the Mo content is preferably set to 0.010% or more. The Mo content may be 0.020% or more, 0.050% or more, or 0.080% or more.

On the other hand, if the Mo content exceeds 1.000%, the effect of suppressing ferrite transformation 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 is an element effective for increasing the strength of steel, as well as Mn, and suppresses pearlite transformation. In addition, cr is also an element that contributes to improvement of grain boundary bonding strength. Therefore, cr may be contained. In order to obtain the above-described effects, the Cr content is preferably set to 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 carbide is formed in the center segregation portion, and the hydrogen embrittlement resistance may be lowered. 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 the improvement of the strength of a steel sheet by precipitate strengthening, grain strengthening by suppressing ferrite grain growth, and dislocation strengthening by suppressing recrystallization. In addition, ti is an element that contributes to improving the grain boundary bonding strength. Therefore, ti may be contained. In order to obtain the above-described effects, the Ti content is preferably set to 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%, precipitation of carbonitrides increases, and hydrogen embrittlement resistance may be deteriorated. 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 and promotes the formation of a low-temperature transformation structure such as bainite or martensite during cooling from the austenite temperature region. In addition, B is an element that is beneficial for increasing the strength of steel. In addition, B is an element that contributes to improvement of grain boundary bonding strength. Thus, B may be contained. In order to obtain the above-mentioned effects, the B content is preferably set to 0.0001% or more. The B content may be 0.0003% or more, 0.0005% or more, or 0.0010% or more.

On the other hand, if the B content exceeds 0.0100%, coarse B oxides are formed in the steel. Since this oxide becomes a starting point of void generation during cold working, the hydrogen embrittlement resistance may be deteriorated by the generation of coarse B oxide. Therefore, the B content is set to 0.0100% or less. The B content may be 0.0080% or less, 0.0060% or less, or 0.0050% or less.

Nb：0～0.500％

Nb is an element effective for controlling the form of carbide like Ti, and is also an element effective for improving toughness by refining the structure. In addition, nb is an element that contributes to improvement of grain boundary bonding strength. Therefore, nb may be contained. In order to obtain the above-described effects, the Nb content is preferably set to 0.001% or more. The Nb content may be 0.002% or more, 0.010% or more, or 0.020% or more.

On the other hand, if the Nb content exceeds 0.500%, the formation of coarse Nb carbides becomes remarkable. Since cracks are likely to occur in the coarse Nb carbide, the hydrogen embrittlement resistance may be deteriorated by the formation of the coarse Nb carbide. Therefore, the Nb content is set to 0.500% or less. The Nb content may be 0.450% or less, 0.400% or less, or 0.300% or less.

V：0～0.500％

V is an element that contributes to the improvement of the strength of the steel sheet by precipitate strengthening, grain strengthening based on the suppression of ferrite grain growth, and dislocation strengthening via the suppression of recrystallization. In addition, V is also an element that contributes to improving the grain boundary bonding strength. And thus may also contain V. In order to obtain the above-mentioned effects, the V content is preferably set to 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%, precipitation of carbonitrides increases, and hydrogen embrittlement resistance may be deteriorated. 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 for improving the strength of the steel sheet. If the Cu content is less than 0.001%, a sufficient effect cannot be obtained. Therefore, in order to obtain the above-described effects, the Cu content is preferably set to 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 be deteriorated. In addition, if the Cu content is high, the steel material is embrittled during hot rolling, and hot rolling may not be performed. Therefore, 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 for improving the strength of the steel sheet. In addition, W forms precipitates and crystals. Since precipitates and crystals containing W serve as hydrogen trapping sites, W is an element effective for improving hydrogen embrittlement resistance. In addition, W is also an element that contributes to improvement of grain boundary bonding strength. Therefore, W may be contained. In order to obtain the above-mentioned effects, the W content is preferably set to 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, when the W content exceeds 0.100%, the formation of coarse W precipitates or crystals becomes remarkable. The coarse W precipitates or crystals are prone to cracking, and the cracks propagate in the steel under low load stress. Therefore, if coarse W precipitates or crystals are formed, hydrogen embrittlement resistance may be deteriorated. Therefore, 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 is an element effective for controlling the form of carbide and improving the strength of steel sheet, similarly to Nb, V and W. In addition, ta is an element that contributes to improvement of grain boundary bonding strength. Thus, ta may also be present. In order to obtain the above-described effects, the Ta content is preferably set to 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 amount of fine Ta carbide precipitates, and as the strength of the steel sheet increases, ductility may decrease or bending resistance and hydrogen embrittlement resistance may decrease. Therefore, 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 capable of controlling the form of sulfide in a trace amount. Mg may also be contained. In order to obtain the above-described effect, the Mg content is preferably set to 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 are formed, and hydrogen embrittlement resistance may be lowered. 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％

In addition to being useful as a deoxidizing element, ca is also an effective element for controlling the morphology of sulfides. Therefore, ca may be contained. In order to obtain the above-mentioned effects, the Ca content is preferably set to 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 be lowered. Therefore, the Ca content was 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 sulfide in a trace amount, similarly to Mg and Ca. And thus may also contain Y. In order to obtain the above-mentioned effects, the Y content is preferably set to 0.001% or more. The Y content may be 0.002% or more, 0.004% or more, or 0.006% or more.

On the other hand, if the Y content exceeds 0.050%, coarse Y oxides are formed, and hydrogen embrittlement resistance may be lowered. 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 form of sulfide in a trace amount, similarly to Mg, ca, and Y. Zr is also an element that contributes to improvement of grain boundary bonding strength. Therefore, zr may be contained. In order to obtain the above-mentioned effects, the Zr content is preferably set to 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 oxide is formed, and hydrogen embrittlement resistance may be lowered. Therefore, the Zr content was 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 morphology of sulfide in a trace amount, as in Mg, ca, Y, zr. Therefore, la may be contained. In order to obtain the above-mentioned effects, the La content is preferably set to 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 oxide is formed, and hydrogen embrittlement resistance may be lowered. Therefore, the La content was 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 sulfide in a trace amount, similar to La. And thus may also contain Ce. In order to obtain the above-described effects, the Ce content is preferably set to 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 oxide is formed, and the hydrogen embrittlement resistance may be lowered. Therefore, the Ce content was 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 lowered due to grain boundary embrittlement. If the Sn content exceeds 0.050, the adverse effect becomes particularly remarkable. Therefore, the Sn content is set to 0.050% or less. The Sn content may be 0.040% or less, 0.030% or less, or 0.020% or less.

The smaller the Sn content, the more preferable the Sn content is, but when the Sn content is set to less than 0.001%, the refining cost increases. Therefore, the Sn content may be set to 0.001% or more. The Sn content may be 0.002% or more, 0.005% or more, or 0.010% or more.

Sb：0～0.050％

Sb is an element contained in the case of using scrap as a steel raw material, similarly to Sn. Sb is an element that strongly segregates to grain boundaries, causing grain boundary embrittlement and a decrease in ductility. If the Sb content exceeds 0.050, the adverse effect thereof becomes particularly remarkable. The Sb content was thus set to 0.050%. The Sb content may be 0.040% or less, 0.030% or less, or 0.020% or less.

The smaller the Sb content, the more preferable the Sb content, but when the Sb content is set to less than 0.001%, the refining cost increases. Therefore, the Sb content may be set to 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％

As is an element which is contained in the case of using scrap As a steel raw material, and which is strongly segregated to grain boundaries, and causes grain boundary embrittlement and a decrease in ductility, like Sn and Sb. If the As content is large, the hydrogen embrittlement resistance may be lowered. If the As content exceeds 0.050, the adverse effect thereof becomes particularly remarkable. The As content was thus set to 0.050%. The As content may be 0.040% or less, 0.030% or less, or 0.020% or less.

The smaller the As content, the more preferable the As content, but when the As content is set to less than 0.001%, refining cost increases. Therefore, the As content may be set to 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 sheet according to the present embodiment may contain the basic component, and the remainder may contain Fe and impurities, or may contain 1 or more of the basic components, and the remainder may contain Fe and impurities.

The chemical composition of the steel sheet according to the present embodiment may be measured by a general method. For example, as long as it is in accordance with JIS G1201:2014, and measuring the cuttings by adopting ICP-AES (inductively coupled plasma atomic emission spectrometry: inductively Coupled Plasma-Atomic Emission Spectrometry). In this case, the chemical composition is the average content of the total plate thickness. The elements C and S which cannot be measured by ICP-AES may be measured by a combustion-infrared absorption method, N may be measured by an inert gas fusion-thermal conductivity method, and O may be measured by an inert gas fusion-non-dispersive infrared absorption method.

When the steel sheet has a coating layer on the surface, chemical composition analysis may be performed after removing the coating layer by mechanical grinding or the like. In the case where the coating layer is a plating layer, the plating layer may be removed by dissolving the plating layer with an acid solution to which a corrosion inhibitor for inhibiting corrosion of the steel sheet is added.

< Microstructure (Metal Structure) >

Next, the microstructure of the steel sheet according to the present embodiment will be described. In the present embodiment, the microstructure is a microstructure at a position (t/4 portion) of 1/8 to 3/8 of the plate thickness from the surface of the steel plate in the plate thickness direction. The microstructure of the t/4 portion is defined because it is a representative microstructure of the steel sheet and is closely related to the properties of the steel sheet.

The fraction (%) of each phase is the area ratio unless otherwise specified.

Ferrite: 5.0% or less

Ferrite exerts an influence on the deformability of the steel having a martensite-based structure. As the area ratio of ferrite increases, the local deformability and hydrogen embrittlement resistance decrease. In particular, if the area ratio of ferrite exceeds 5.0%, the hydrogen embrittlement resistance may be lowered due to fracture in elastic deformation at the time of stress loading. Therefore, the area ratio of ferrite is set to 5.0% or less. The area ratio of ferrite may be 4.0% or less, 3.0% or less, or 2.0% or less.

The area ratio of ferrite may be 0%, but if it is set to less than 1.0%, strict control is required in production, resulting in a decrease in yield. Therefore, the area ratio of ferrite may be set to 1.0% or more.

Martensite and tempered martensite: total more than 90.0%

The total area ratio of martensite and tempered martensite affects the strength of the steel, and the tensile strength increases as the area ratio increases. If the total area ratio of martensite and tempered martensite is 90.0% or less, the targeted tensile strength cannot be achieved. Further, the fracture may be caused during elastic deformation under a stress load, and the non-uniformity of the microstructure may be increased by the formation of a structure other than martensite and tempered martensite, thereby reducing the hydrogen embrittlement resistance. Therefore, the total area ratio of martensite and tempered martensite is set to be more than 90.0%. The total area ratio of the martensite and tempered martensite may be 95.0% or more, 97.0% or more, 99.0% or more, or 100.0%.

The remainder: comprises 1 or more than 2 of bainite, pearlite and retained austenite

The area ratio of the structure (the remaining structure) other than the above structure may be 0%, but when the remaining structure is present, the remaining structure contains 1 or 2 or more of bainite, pearlite, and retained austenite.

If the area ratio of the remaining part of the structure exceeds 8.0%, fracture in elastic deformation at the time of stress load may occur, and hydrogen embrittlement resistance may be lowered. Therefore, the area ratio of the remaining tissue is preferably 8.0% or less, more preferably 7.0% or less. Among them, pearlite and retained austenite are particularly preferable because they are structures that deteriorate the local ductility of steel.

On the other hand, if the area ratio of the remaining structure is set to 0%, strict control is required in manufacturing, and thus the yield may be lowered. Therefore, the area ratio of the remaining tissue may be 1.0% or more.

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

(Method for evaluating area ratio of ferrite)

The area ratio of ferrite can be obtained by observing the t/4 portion (the range of 1/8 to 3/8 of the plate thickness centered on a position 1/4 of the plate thickness from the surface in the plate thickness direction) by using an electron channel contrast image of a Field Emission scanning electron microscope (FE-SEM: field Emission-Scanning Electron Microscope). The electron channel contrast image is a method of detecting a difference in crystal orientation within a crystal grain as a difference in contrast of the image, and in the image, a portion of a structure determined to be not pearlite, bainite, martensite, retained austenite, but ferrite is reflected as polygonal ferrite in a uniform contrast. For 8 fields of the 35 μm×25 μm electron channel contrast image, the area ratio of polygonal ferrite in each field was calculated by image analysis, and the average value was used as the area ratio of ferrite.

(Method for evaluating the area ratio of martensite and tempered martensite in total)

The martensite and tempered martensite may be obtained as a total area ratio from the image captured by the electron channel contrast. These structures are less likely to be corroded than ferrite, and therefore exist as projections on the structure observation surface. Tempered martensite is an aggregation of lath-shaped grains, and contains iron-based carbides having a long diameter of 20nm or more, which belong to a plurality of varieties, i.e., a plurality of iron-based carbide groups extending in different directions, inside. The retained austenite also exists as a convex portion on the tissue observation surface. Therefore, the area ratio of the sum of the martensite and tempered martensite can be accurately measured by subtracting the area ratio of the convex portion obtained by the above-described steps from the area ratio of the retained austenite measured by the steps described later.

(Method for evaluating the area ratio of the total of bainite, pearlite, and retained austenite)

The area ratio of the retained austenite can be calculated by measurement using X-rays. That is, the polishing solution was removed from the plate surface of the sample to a position 1/4 of the plate thickness in the plate thickness direction by mechanical polishing and chemical polishing. Then, for the sample after grinding, the structure fraction of the retained austenite was calculated from the integrated intensity ratios of 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, and was taken as the area fraction of the retained austenite.

The pearlite can be obtained from an image captured by the electron channel contrast as described above. Pearlite is a structure in which plate-like carbide and ferrite are arranged.

The bainite is a group of lath-shaped grains, and contains no iron-based carbide having a length of 20nm or more in the interior, or contains iron-based carbide having a length of 20nm or more in the interior, and the carbide belongs to a single variety, i.e., an iron-based carbide group extending in the same direction. The term "iron-based carbide group extending in the same direction" means that the difference in the extending direction of the iron-based carbide group is within 5 °.

< Binding Strength energy >

Cracks due to hydrogen embrittlement occur because hydrogen in steel segregates to grain boundaries to reduce the bonding strength of the grain boundaries, and the grain boundaries with reduced bonding strength become the starting points of the cracks. In contrast, if a predetermined alloy element is segregated in the grain boundary, the bonding strength of the grain boundary is improved. In addition, the hydrogen that has intruded is difficult to segregate to the grain boundaries by the segregation of the alloy element, and even if hydrogen intrudes, the decrease in the grain boundary bonding strength due to hydrogen can be suppressed.

In the microstructure mainly composed of martensite and tempered martensite, the inventors examined the contribution of each alloy element to the improvement of the grain boundary strength by using as a main grain boundary a prior austenite grain boundary which is an interface (interface of martensite/tempered martensite, martensite/martensite, or tempered martensite/tempered martensite) in which the difference in orientation between adjacent martensite and tempered martensite is 15 ° or more. As a result, it was found that the bonding strength energy E _GB can be expressed as shown in the following formula (1) using the concentration of each alloy element at the grain boundary; when E _GB is 0.50 or more, the hydrogen embrittlement resistance is remarkably improved.

Therefore, in the steel sheet according to the present embodiment, when the interface where the difference in orientation between the adjacent martensite and tempered martensite is 15 ° or more is defined as the prior austenite grain boundary, the bonding strength energy E _GB determined by the concentration of each alloy element in the prior austenite grain boundary satisfies the following formula (1).

Wherein the chemical symbol in the formula represents the concentration of each alloy element in mass% on the prior austenite grain boundary.

The reason why the interface with the azimuth difference of 15 ° or more is aimed is that hydrogen tends to accumulate preferentially in the prior austenite grain boundaries of 15 ° or more.

As is clear from the formula (1), not all the alloy elements segregated to the grain boundaries can improve the bonding strength energy, and the grain boundary bonding energy can be improved by segregating a large amount of the alloy elements capable of improving the grain boundary bonding energy.

The term "boundary surface between adjacent martensite and tempered martensite having a difference in orientation of 15 ° or more" means a boundary surface between martensite and martensite having a difference in orientation of 15 ° or more, a boundary surface between martensite and tempered martensite having a difference in orientation of 15 ° or more, and a boundary surface between tempered martensite and tempered martensite having a difference in orientation of 15 ° or more.

The concentration of each alloy element in the prior austenite grain boundary can be determined as follows: EDS (energy dispersive X-ray spectrometer) using TEM (transmission electron microscope) was observed at the t/4 section (in the range of 1/8 to 3/8 of the thickness of the sheet centered at a position 1/4 of the thickness of the sheet from the surface in the sheet thickness direction) in the same manner as in the SEM observation described above. More specifically, as a TEM (transmission electron microscope), a TEM transmission electron microscope (Cs-corrected TEM) with spherical aberration correction is used.

The sheet sample used for TEM observation can be obtained by the following method. Samples for measuring the amount of the alloy element were collected from the steel sheet samples from the range of 1/8 to 3/8 of the plate thickness, and the samples were polished to a thickness of about 100 μm by wet polishing using a coated diamond abrasive.

Then, electrolytic grinding was performed by electrolytic double jet grinding to a thinness that can be observed by TEM. The electrolytic grinding method is carried out by adopting an electrolytic double-jet grinding device. Since the appropriate conditions for electrolytic double jet polishing vary depending on the base material composition of the sample, it is necessary to sample each sample. After the double spraying, the thin sheet sample was uniformly polished by Ar ion polishing, and the accuracy of quantification of the element on the prior austenite grain boundary was improved.

The thus obtained sheet sample was observed by Cs-corrected TEM. The observation position was on the prior austenite grain boundaries, and the prior austenite grain boundaries were found in the following manner. In Bright-field images (Bright-FIELD IMAGE) in TEM observation, prior austenite grain boundaries, lath bundle (packet) boundaries, and lath block (block) boundaries appear as black lines when observed at 30000 times. The sample is tilted and rotated so that the black line representing any prior austenite grain boundary among these black lines is horizontal to the incidence direction of the electron beam of the TEM. In this state, elemental analysis by EDS was performed 100000 times directly above the prior austenite grain boundaries. Elemental analysis using the cumulative number of EDS analysis was performed as follows. And 3 times of point analysis are carried out right above the prior austenite grain boundary, so that the concentration of alloy elements on the prior austenite grain boundary is quantified. This analysis was performed on 5 prior austenite grain boundaries, and the average alloy element concentration was calculated. The average concentration of the alloying element was taken as the concentration of the alloying element at the prior austenite grain boundary.

(Mechanical Properties)

In the steel sheet according to the present embodiment, the Tensile Strength (TS) is set to 1500MPa or more as a strength contributing to weight saving of the automobile body.

The upper limit is not necessarily limited, but if the tensile strength is increased, the formability may be lowered, and therefore the tensile strength may be set to 2000MPa or less.

(Plate thickness)

The thickness of the steel sheet according to the present embodiment is not limited, but is preferably 1.0 to 2.2mm. The thickness is more preferably 1.05mm or more, and still more preferably 1.1mm or more. The thickness is more preferably 2.1mm or less, and still more preferably 2.0mm 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. The coating layer may be formed of zinc, aluminum, magnesium, or an alloy and impurities thereof.

The surface is provided with a coating layer to improve corrosion resistance. If there is a concern about pits caused by corrosion, the steel sheet for automobiles may not be thinned to a certain fixed plate thickness or less even if the steel sheet is made stronger. One of the purposes of increasing the strength of steel sheets is to reduce the weight due to the reduction in thickness, and therefore, even if high-strength steel sheets are developed, the application sites thereof are limited if the corrosion resistance is low. As a method for solving these problems, forming a coating layer on both front and back surfaces can be considered in order to improve corrosion resistance.

Even if a coating layer is formed, the hydrogen embrittlement resistance of the steel sheet according to the present embodiment is not impaired.

Examples of the coating layer include a hot dip galvanized layer, an alloyed hot dip galvanized layer, an electro-galvanized layer, an aluminized layer, a Zn-Al alloy coating layer, an Al-Mg alloy coating layer, and a Zn-Al-Mg alloy coating layer.

When the coating layer is provided on the surface, the surface to be the reference of the t/4 section is the surface from which the base metal of the coating layer is removed.

< Manufacturing method >

The steel sheet according to the present embodiment can be manufactured by a manufacturing method including the following steps (I) to (VII), as long as the steel sheet has the above-described characteristics, regardless of the manufacturing method.

(I) A heating step of heating a billet having a predetermined chemical composition,

(II) a hot rolling step of hot-rolling the heated slab to obtain a hot-rolled steel sheet,

(III) a cooling step of cooling the hot-rolled steel sheet to a coiling temperature of 550 to 700 ℃ at an average cooling rate of 20 to 50 ℃ per second, starting cooling within 3.0 seconds from the end of the hot-rolling step,

(IV) a coiling step of coiling the hot-rolled steel sheet after the cooling step at the coiling temperature,

(V) a holding step of holding the hot-rolled steel sheet after the coiling step in a temperature range of 400 to 550 ℃ for 600 seconds or longer,

(VI) a cold rolling step of pickling and cold-rolling the hot-rolled steel sheet after the dwell step to obtain a cold-rolled steel sheet,

(VII) an annealing step of annealing the cold-rolled steel sheet after the cold-rolling step by maintaining the cold-rolled steel sheet at an annealing temperature of 800 ℃ or higher and lower than 900 ℃.

Hereinafter, preferable conditions in each step will be described.

(Heating step)

In the heating step, a billet such as a slab having the same chemical composition as the steel sheet according to the present embodiment is heated before 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 ℃.

The billet to be used is preferably cast by continuous casting from the viewpoint of productivity, but may be produced by ingot casting or thin slab casting.

In the case where a billet obtained by continuous casting can be supplied to the hot rolling step in an original state at a sufficiently high temperature, the heating step may be omitted.

(Hot Rolling Process)

In the hot rolling step, the heated slab is hot-rolled to obtain a hot-rolled steel sheet.

The hot rolling step includes rough rolling and finish rolling, wherein the finish rolling is performed in multiple passes, 4 or more passes among the multiple passes are set as large reduction passes with a reduction ratio of 20% or more, and the inter-pass time of each of the large reduction passes is set to 5.0 seconds. The rolling start temperature is set to 950 to 1100 ℃ and the rolling end temperature is set to 800 to 950 ℃.

In finish rolling, a large reduction pass with a reduction ratio of 20% or more: 4 or more passes ]

Inter-pass time: within 5.0 seconds ]

By controlling the reduction rate, the number of rolling passes, and the time between passes in finish rolling, the morphology of austenite grains can be controlled to be equiaxed and fine. If the austenite grains are equiaxed and fine, the grain boundary diffusion of the alloy element is promoted, and the precipitation of alloy carbide or nitride is promoted at the grain boundary. If the pass of the reduction ratio of 20% or more (large reduction pass) is less than 4 passes, sufficient effect cannot be obtained due to the austenite remaining without recrystallization. Therefore, the reduction ratio is set to 20% or more in 4 or more passes (4 or more passes of reduction is performed at a reduction ratio of 20% or more). The reduction ratio is preferably set to 20% or more in 5 or more passes. On the other hand, the upper limit of the number of passes of the reduction ratio of 20% or more is not particularly limited, but if the number of passes exceeds 10 passes, a plurality of rolling stands are required, which may lead to an increase in the size of the equipment and an increase in the manufacturing cost. Therefore, the number of passes (number of passes) of the reduction ratio of 20% or more may be 10 passes or less, 9 passes or less, or 7 passes or less.

In addition, the inter-pass time in finish rolling exerts a large influence on recrystallization and grain growth of austenite grains after rolling. Even when the number of large reduction passes is set to 4 or more, if the inter-pass time exceeds 5.0 seconds, grain growth tends to occur, and austenite grains coarsen.

On the other hand, the lower limit of the inter-pass time is not limited, but if the inter-pass time of each large reduction pass is less than 0.2 seconds, the recrystallization of austenite is not completed, and the proportion of unrecrystallized austenite increases, and a sufficient effect may not be obtained. Therefore, the inter-pass time of the large reduction pass is preferably set to 0.2 seconds or more. The inter-pass time may be 0.3 seconds or more or 0.5 seconds or more.

(Cooling step)

In the cooling step, cooling is started within 3.0 seconds from the end of the hot rolling step (the end of the final pass of the finish rolling), and the hot rolled steel sheet after the hot rolling step is cooled to a coiling temperature of 550 to 700 ℃ at an average cooling rate of 20 to 50 ℃/sec.

When 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 ℃/sec, ferrite transformation is generated from austenite before coiling. In this case, the driving force of the precipitate decreases, and it is difficult for the precipitate to be uniformly and finely precipitated in the subsequent step.

On the other hand, when the average cooling rate to the winding temperature is too high, a hardened phase becomes easy to form. In this case, the manufacturability thereafter is significantly deteriorated, and the productivity is lowered.

If the cooling stop temperature is less than 550 ℃, the precipitation of the precipitate is delayed, and the manufacturability is deteriorated, resulting in a decrease in productivity. Further, if the cooling stop temperature exceeds 700 ℃, ferrite transformation occurs from austenite, and the driving force for precipitation of carbide is reduced, so that it is difficult for the precipitates to be uniformly and finely precipitated in the subsequent step, which is not preferable. Further, if the cooling stop temperature exceeds 700 ℃, an internal oxide layer tends to be formed on the surface of the steel sheet, and cracks tend to occur on the surface, or productivity in acid washing in the subsequent steps is significantly deteriorated, which is not preferable.

The lower limit of the time from the end of hot rolling to the start of rolling is not limited, and the rolling may be performed in a short time as much as possible within the range of facility restrictions.

(Winding Process)

(Residence Process)

In the coiling step, the hot-rolled steel sheet after the cooling step is coiled at a coiling temperature (cooling stop temperature). In the subsequent holding step, the hot-rolled steel sheet after coiling is held (held) at a temperature range of 400 to 550 ℃ for 600 seconds or longer. By controlling the coiling and stopping conditions, alloy carbide or nitride can be precipitated in the steel sheet.

The precipitates deposited here can be unevenly distributed in the prior austenite grain boundaries by controlling the subsequent steps.

If the holding temperature is too high, the precipitates coarsen and are not uniformly dispersed. In addition, if the holding temperature is too low, although the precipitates are miniaturized, a lot of time is required until the end of precipitation, and manufacturability and productivity are lowered. Further, if the holding time is short, alloy carbide is not sufficiently precipitated.

In the case of residence under the above conditions, for example, a capping cover, a coating with a heating box or the like may be used.

(Cold Rolling Process)

In the cold rolling step, the hot-rolled steel sheet after the staying step is unwound, and subjected to pickling and cold rolling to obtain a cold-rolled steel sheet.

The scale on the surface of the hot-rolled steel sheet is removed by pickling, whereby the chemical conversion treatability and the plating property of the cold-rolled steel sheet can be improved. The acid washing may be performed under known conditions, and may be performed once or in a plurality of times. The rolling reduction (rolling reduction) of the cold rolling is not particularly limited. For example, 20 to 80%.

(Annealing step)

In the annealing step, the cold-rolled steel sheet after the cold-rolling step is annealed while being held at an annealing temperature of 800 ℃ or higher and lower than 900 ℃.

In this annealing step, these precipitates act as pinning sites to the prior austenite grain boundaries in a relatively low temperature region during heating to the austenite single-phase region, i.e., the annealing temperature. As a result, the precipitates are unevenly distributed on the prior austenite grain boundaries. If the heating is further carried out to a relatively high temperature region, the precipitate melts due to thermal instability. As a result, the alloy element can be segregated in the prior austenite grain boundary.

In this state, as in the post-annealing cooling step described later, austenite is transformed into martensite by quenching, and a high-strength steel sheet in which alloy elements are unevenly distributed in the prior austenite grain boundaries can be obtained.

If the annealing temperature is lower than 800 ℃, the amount of austenite formed is small, and carbide is not sufficiently melted. Therefore, the annealing temperature was set to 800℃or higher. The annealing temperature is preferably 830℃or higher.

On the other hand, if the annealing temperature is 900 ℃ or higher, grain growth occurs, and the prior austenite grain diameter coarsens, so that segregation of a predetermined alloy element to grain boundaries is suppressed, and hydrogen embrittlement resistance may be deteriorated.

The holding time at the annealing temperature is not limited, but if the holding time is less than 10 seconds, there is a concern that the fraction of austenite at the annealing temperature is insufficient or that the melting of carbide is insufficient. Therefore, the holding time is preferably 10 seconds or longer. On the other hand, even if the holding time is prolonged, there is no problem in characteristics, but if continuous annealing is employed, the line length of the equipment is prolonged, and therefore, about 600 seconds may be taken as a substantial upper limit.

The average heating rate to the annealing temperature is preferably 2 to 35 ℃/sec.

(Post-annealing Cooling step)

(Tempering step)

In the post-annealing cooling step, the cold-rolled steel sheet may be cooled to 25 to 300 ℃ at an average cooling rate of 20 to 100 ℃/sec from the annealing temperature after the annealing step.

By this cooling, the steel sheet is quenched in a state where the alloy element is segregated to the austenite grain boundaries, and austenite is transformed into martensite. As a result, a steel sheet having a structure mainly composed of martensite, in which the alloy segregates to the prior austenite grain boundaries, can be obtained.

If the average cooling rate is less than 20 deg.c/sec, a sufficient amount of martensite is not generated. On the other hand, when the average cooling rate exceeds 100 ℃/sec, if continuous annealing is assumed, the capacity of the apparatus is insufficient, and the apparatus may need to be enhanced, so that 100 ℃/sec is a substantial upper limit.

If the cooling stop temperature exceeds 300 ℃, the non-phase-transformed austenite that does not undergo martensitic transformation is liable to undergo bainitic transformation, and the strength may be lowered. On the other hand, even if the cooling stop temperature is set to 25 ℃, the effect is saturated, and the manufacturability is reduced or the cost is increased due to the need for a special cooling medium or the like.

The cold-rolled steel sheet after the post-annealing cooling step may be further subjected to a tempering step of heating to 50 to 550 ℃ and holding for 10 to 1000 seconds. By performing this tempering, the alloy elements that are not completely segregated to the grain boundaries and are present in the grains can be segregated to the prior austenite grain boundaries. In addition, by making the martensite into tempered martensite, formability can be improved.

If the tempering temperature (holding temperature) is lower than 50 ℃ or the holding time is lower than 10 seconds, the above-mentioned effect is not obtained. On the other hand, if the tempering temperature exceeds 550 ℃, there is a case where a decrease in strength due to a decrease in dislocation density in tempered martensite occurs, resulting in a decrease in tensile strength. In addition, carbides are coarsely precipitated on the prior austenite grain boundaries, and the hydrogen embrittlement resistance may be deteriorated. Further, if the holding time exceeds 1000 seconds, the strength is lowered and the productivity is lowered. Tempering may be performed in a continuous annealing facility, or may be performed off-line after continuous annealing by another facility.

In the post-annealing cooling step, the temperature range of 350 to 650 ℃ may be maintained during cooling (the 2 nd temperature range: a temperature range which is considered to be higher than the Ms point) for 10 to 200 seconds. In this case, the cooling rate up to the 2 nd temperature region and the average cooling rate from the holding temperature to 25 to 300 ℃ (cooling stop temperature) may be set to 20 to 100 ℃/sec, respectively, in addition to the holding.

That is, in this case, after the annealing step, the annealing temperature is cooled at an average cooling rate of 20 to 100 ℃/sec until reaching a temperature range of 350 to 650 ℃, the holding is performed in the temperature range of 2 for 10 to 200 seconds, and the annealing temperature is cooled at an average cooling rate of 20 to 100 ℃/sec until reaching 25 to 300 ℃.

By holding at 350 to 650 ℃, alloy elements existing in the crystal grains which are not completely segregated to the grain boundaries can be segregated to the prior austenite grain boundaries, and the hydrogen embrittlement resistance can be improved. However, if the holding temperature is lower than 350 ℃, bainite transformation tends to occur, and there is a concern that the strength may be lowered. If the holding time is less than 20 seconds, the effect of segregating the elements present in the grains to the prior austenite grain boundaries is not obtained. On the other hand, if the holding temperature exceeds 650 ℃, ferrite transformation from austenite becomes easy to occur, resulting in a decrease in tensile strength. Further, if the holding time exceeds 200 seconds, bainitic transformation or ferrite transformation from austenite becomes easy to occur. The holding temperature is preferably in the range of 370 ℃ to 630 ℃, more preferably 390 ℃ to 610 ℃. The holding time is preferably 30 seconds to 180 seconds, more preferably 50 seconds to 160 seconds.

The holding and tempering steps in the middle of the cooling step after annealing further promote the segregation of the alloy element to the prior austenite grain boundary. Therefore, either one of the steps may be performed, or both of the steps may be performed. Either step may not be performed.

In the method for producing a steel sheet according to the present embodiment, a coating layer forming step of forming a coating layer on a surface (one surface or both surfaces) of the steel sheet may be provided.

As the coating layer, a coating layer containing zinc, aluminum, magnesium, or an alloy thereof is preferable. The coating layer is, for example, a plating layer.

The coating method is not limited, and for example, when a coating layer mainly composed of zinc is formed by hot dip plating, a condition in which a cold-rolled steel sheet is adjusted (by heating or cooling) until the temperature of the steel sheet reaches (bath temperature-40) to (bath temperature +50) c, and then immersed in a bath at 450 to 490 c to form a coating layer can be exemplified.

This condition is preferable because if the steel sheet temperature at the time of immersion in the plating bath is lower than the hot dip galvanization bath temperature of-40 ℃, the heat release at the time of immersion in the plating bath increases, and a part of the hot dip galvanization solidifies, so that the appearance of the plated layer may be deteriorated, and if the hot dip galvanization bath temperature exceeds +50℃, there is a possibility that an operational problem is induced that the plating bath temperature increases.

When forming a zinc-based plating layer, the composition of the plating bath is preferably: the effective Al content (the value obtained by subtracting the total Fe content from the total Al content in the plating bath) is 0.050 to 0.250 mass%, and if necessary, mg is contained, and the balance is Zn and impurities. If the effective Al amount in the plating bath is less than 0.050 mass%, the penetration of Fe into the plating layer is excessively performed, and there is a concern that the plating layer adhesion is lowered. On the other hand, if the effective Al amount in the plating bath exceeds 0.250 mass%, al-based oxides that inhibit movement of Fe atoms and Zn atoms may be generated at the boundary between the steel sheet and the plating layer, and the adhesion of the plating layer may be lowered.

The formation of the coating layer may be performed after the annealing and cooling step, or may be performed in the annealing and cooling step or the tempering step. That is, the cooling step may be performed as a part of holding at 350 to 650 ℃ in the post-annealing cooling step or as a part of holding at 50 to 550 ℃ in the tempering step.

When a plating layer mainly composed of zinc is formed as the coating layer, an alloying treatment may be further performed. In this case, the condition of holding the steel sheet on which the plating layer is formed at 480 to 550 ℃ for 1 to 30 seconds is exemplified.

The alloying step may be performed in the annealing-after-cooling step or the tempering step. That is, the cooling step may be performed as a part of holding at 350 to 650 ℃ in the post-annealing cooling step or as a part of holding at 50 to 550 ℃ in the tempering step.

The surface of the coating layer may be subjected to upper layer plating, various treatments such as chromate treatment, phosphate treatment, treatment for improving lubricity, and treatment for improving weldability, for the purpose of improving coatability and weldability.

Examples

Hereinafter, embodiments of the present invention are shown. The embodiment shown below is an example of the present invention, and the present invention is not limited to the embodiment described below.

Example 1]

Steels having chemical compositions shown in tables 1-1 to 1-4 were melted and cast into billets.

The slab was inserted into a furnace heated to 1220 ℃, left for 60 minutes, taken out in the atmosphere, and hot-rolled to obtain a steel sheet (hot-rolled steel sheet) having a sheet thickness of 2.8 mm. In the hot rolling, a rolling mill having 7 rolling stands was used, and all of the finish rolling was performed 7 times continuously (in a manner of fixing the inter-pass time), wherein the rolling passes having a reduction ratio exceeding 20% were performed 4 times. In finish rolling, the inter-pass time between each rolling pass to which a reduction of 20% or more is applied and the rolling pass before 1 pass among the rolling passes is set to 0.6 seconds. The finish rolling was started at 1060℃and ended at 870 ℃.

After the completion of hot rolling, the hot rolled steel sheet was cooled by water cooling after 2.2 seconds, cooled to 580 ℃ at an average cooling rate of 38.0 ℃/sec, and charged into a furnace at 530 ℃ after coiling, and held for 1800 seconds.

Next, the scale on the hot-rolled steel sheet was removed by pickling, and then cold rolling was performed with a reduction of 50.0%, thereby obtaining a cold-rolled steel sheet having a sheet thickness of 1.4 mm.

The cold-rolled steel sheet was heated to 880℃at an average heating rate of 12.0℃per second, held at 880℃for 120 seconds, and then cooled to 150℃at an average cooling rate of 42.0℃per second.

Then, tempering was performed in which the cold-rolled steel sheet was reheated to 230 ℃ and maintained for 180 seconds. No plating treatment was performed.

The chemical composition was analyzed by a sample collected from the resulting steel sheet. As a result, the chemical compositions were the same as those of the steels shown in tables 1-1 to 1-4.

The area ratio of ferrite, martensite, tempered martensite, and the remainder (1 or 2 or more of bainite, pearlite, and retained austenite) in the microstructure of the t/4 portion was obtained for the obtained cold-rolled steel sheet by the above-described method.

The results are shown in Table 2.

Further, the concentration of each alloy element at the austenite grain boundary was measured by the above method, and E _GB was obtained.

The obtained cold-rolled steel sheet was evaluated for tensile strength, total elongation, and hydrogen embrittlement resistance (hydrogen embrittlement resistance) in the following manner.

(Evaluation method of tensile Property)

The tensile test was performed by collecting a test piece of JIS No. 5 from a direction in which the longitudinal direction of the test piece was parallel to the rolling direction of the steel strip in accordance with JIS Z2241 (2011), and measuring the Tensile Strength (TS) and the total elongation (El).

(Evaluation method of Hydrogen embrittlement resistance)

The hot dip galvanized steel sheet manufactured by the method for manufacturing a steel sheet according to the embodiment of the present invention was evaluated for hydrogen embrittlement resistance by the following method. Specifically, after the steel sheet was cut at 15% clearance, a U-bend 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 fastened with bolts, thereby applying stress. The applied stress may be calculated from the strain of the strain gauge being monitored. The load stress applies a stress corresponding to 80% of the Tensile Strength (TS) (e.g., at a of table 2, the applied stress=1515 mpa×0.8=1212 MPa). This is because it is considered that the residual stress introduced during forming corresponds to the tensile strength of the steel sheet.

The obtained U-bend test piece was immersed in an aqueous HCl solution at 25 ℃ and pH2, and kept for 96 hours, and the presence or absence of cracks was examined. The lower the pH of the aqueous HCl solution and the longer the immersion time, the greater the amount of hydrogen that intrudes into the steel sheet, and therefore the hydrogen embrittlement environment is a severe condition. After dipping, the U-bend test piece was evaluated, and the case where a crack exceeding 1.00mm in length was observed was evaluated as NG, the case where a crack exceeding 1.00mm in length was not observed was evaluated as OK, and the case where OK was evaluated as OK was accepted, and the case where NG was rejected.

A steel sheet having a tensile strength of 1500MPa or more and excellent hydrogen embrittlement resistance was evaluated as a steel sheet having high strength and excellent hydrogen embrittlement resistance, when the evaluation was OK.

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TABLE 2

As can be seen from tables 1 to 2, the production No. P has a low C content, and thus has a tensile strength of less than 1500MPa.

Production No. Q has a high C content, and therefore, decreases hydrogen embrittlement resistance.

Production No. R is reduced in hydrogen embrittlement resistance because the Si content is high, thereby suppressing precipitation of alloy carbide in the hot rolling step and segregation of grain boundary strengthening element to prior austenite grain boundary in the annealing step.

Production No. S has a tensile strength of less than 1500MPa due to its low Mn content.

Production No. T has a high Mn content and a decrease in E _GB, and thus decreases hydrogen embrittlement resistance.

Since the production No. U has a high P content, E _GB is lowered and hydrogen embrittlement resistance is lowered.

Production No. V has a high S content, and therefore, decreases hydrogen embrittlement resistance.

Since the production No. W has a high N content, coarse nitrides are produced, and the hydrogen embrittlement resistance is lowered.

Since the production No. X has a high Al content, coarse Al oxides are formed, and the hydrogen embrittlement resistance is lowered.

Since the production No. Y has a high B content, coarse B oxides are produced, and the hydrogen embrittlement resistance is lowered.

Since the production No. Z has a high Ti content, coarse carbonitrides are formed, and the hydrogen embrittlement resistance is lowered. Further, since coarse carbonitrides are formed, ti segregated in grain boundaries is reduced, and E _GB is lowered. In addition, since the amount of C effective for improving the strength is reduced, the tensile strength is set to 1500MPa or less.

Since the production No. AA has a high Nb content, coarse Nb carbides are formed, and the hydrogen embrittlement resistance is lowered.

Since the production No. AB has a high V content, coarse V carbide is formed, and hydrogen embrittlement resistance is lowered.

The production No. AC was high in O content, and thus formed oxides, which reduced hydrogen embrittlement resistance.

Since the production No. AD has a high Mo content, the precipitation of carbonitrides increases, and the hydrogen embrittlement resistance decreases. In addition, since coarse carbonitrides are formed, mo segregated in grain boundaries is reduced, and E _GB is lowered.

Since the production No. ae has a high Cr content, coarse Cr carbides are formed at the center segregation site in the steel, and the hydrogen embrittlement resistance is reduced. In addition, since coarse carbonitrides are formed, cr segregated to grain boundaries is reduced, and E _GB is lowered.

Production No. AF produced coarse Co carbides due to the high Co content, and reduced hydrogen embrittlement resistance. In addition, coarse carbonitrides are formed, so that Co segregated into grain boundaries is reduced, and E _GB is lowered.

Production No. AG was reduced in hydrogen embrittlement resistance due to the high Ni content.

Since the production No. AH has a high Cu content, E _GB is reduced to less than 0.50, and hydrogen embrittlement resistance is reduced.

Since the production No. AI has a high W content, coarse W precipitates are formed, and the hydrogen embrittlement resistance is reduced. In addition, since the amount of W effective for grain boundary strengthening is reduced, E _GB is less than 0.50, and thus hydrogen embrittlement resistance is reduced.

Since the production No. AJ has a high Ta content, a large amount of fine Ta carbides are precipitated, and the hydrogen embrittlement resistance is reduced.

Since the production No. AK has a high Sn content, E _GB is less than 0.50 by grain boundary embrittlement, and hydrogen embrittlement resistance is reduced.

Since the contents of Sb and As are high in production nos. al and AM, respectively, E _GB is lower than 0.50 by grain boundary segregation, and hydrogen embrittlement resistance is reduced.

Since production No. AN and AO have high Mg and Ca contents, coarse inclusions are formed, and hydrogen embrittlement resistance is reduced.

Production of Nos. AP to AS causes the formation of coarse oxides and decreases hydrogen embrittlement resistance because the contents of Y, zr, la and Ce are high.

In contrast, in production of No. a to O, by appropriately controlling the chemical composition and structure of the steel sheet and the grain boundary strength E _GB of the prior austenite grain boundary, a steel sheet having high strength and excellent hydrogen embrittlement resistance can be obtained.

Example 2]

In order to examine the influence of the production conditions, in example 1, steel grades (steels No. a to O) having excellent characteristics were used, and in the same equipment as in example 1, a steel slab was inserted into a furnace heated to 1250 to 1100 ℃, and after holding for 60 minutes, the steel slab was taken out in the atmosphere, and hot-rolled steel sheets having a sheet thickness of 2.3mm were produced under the production conditions described in tables 3-1 to 3-2. The conditions after coiling are shown in tables 3-1 to 3-4, and cold-rolled steel sheets were obtained under these conditions. A part of the cold-rolled steel sheet was used as a plated steel sheet on which a plating layer was formed. Wherein, symbols GI and GA of the plating treatment represent a method of the galvanization treatment, GI is a steel sheet in which a zinc plating layer is formed on the surface of the steel sheet by immersing the steel sheet in a hot dip galvanization bath at 460 ℃, and GA is a steel sheet in which an alloy layer of iron and zinc is formed on the surface of the steel sheet by heating the steel sheet to 485 ℃ after immersing the steel sheet in the hot dip galvanization bath. In the case of the plating, when the intermediate holding is performed in the 2 nd temperature region, the plating is performed directly (without cooling to room temperature once) after the intermediate holding. When the cooling is not performed in the temperature range 2, the cooling is performed in the middle of cooling to 25 to 300 ℃. In tables 3 to 4, the examples in which tempering was described as "-" are examples in which tempering was not performed. The inter-pass time in the table is the inter-pass time of each pass having a reduction ratio of 20% or more (since rolling is performed by a tandem rolling mill, the inter-pass time is equal to each other). In tables 3 to 3 and 3 to 4, the holding time of the post-annealing cooling step is the holding time in the 2 nd temperature range when cooling to the 2 nd temperature range, but is the holding time in the vicinity of the cooling stop temperature when the cooling stop temperature is outside the 2 nd temperature range.

The area ratio of ferrite, martensite, tempered martensite, and the remainder (1 or 2 or more of bainite, pearlite, and retained austenite) in the microstructure was obtained in the same manner as in example 1, and the concentration of each alloy element in the austenite grain boundary was measured to obtain E _GB.

The obtained cold-rolled steel sheet was evaluated for tensile strength and total elongation in the same manner as in example 1.

The hydrogen embrittlement resistance was evaluated in the following manner.

(Evaluation method of Hydrogen embrittlement resistance)

The hot dip galvanized steel sheet manufactured by the method for manufacturing a steel sheet according to the embodiment of the present invention was evaluated for hydrogen embrittlement resistance by the following method. Specifically, after the steel sheet was cut at 15% clearance, a U-bend 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 fastened with bolts, thereby applying stress. The applied stress may be calculated from the strain of the strain gauge being monitored. The load stress applies a stress corresponding to 80% of the Tensile Strength (TS) (e.g., at a-1 of table 4, the applied stress=1540 mpa×0.8=1232 MPa). This is because it is considered that the residual stress introduced during forming corresponds to the tensile strength of the steel sheet.

The obtained U-bend test piece was immersed in an aqueous HCl solution at 25 ℃ and pH2, and kept for 96 hours, and the presence or absence of cracks was examined. The lower the pH of the aqueous HCl solution and the longer the immersion time, the greater the amount of hydrogen that intrudes into the steel sheet, and therefore the hydrogen embrittlement environment is a severe condition. After dipping, the total length of the cracks of the U-bend test piece (when a plurality of cracks can be seen, as the sum of the values measured one by one) was measured.

The smaller the total length of the crack, the more excellent the hydrogen embrittlement resistance was, and in particular, the case where a crack exceeding 1.00mm in length was observed was evaluated as NG, the case where no crack was observed, and the case where a slight crack having a crack length of 1.00mm or less was observed were evaluated as OK, and the case where no crack was observed and the case where a crack having a crack length of 0.70mm or less were evaluated as Ex. The cases evaluated as OK and Ex were accepted, and the cases evaluated as NG were rejected.

The results are shown in Table 4.

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TABLE 4 Table 4

As can be seen from Table 4, production No. A-2 was carried out for a long period of time from the completion of finish rolling to the start of cooling. Therefore, ferrite transformation during cooling after finish rolling is suppressed, coarsening of pearlite structure is induced, precipitation of alloy carbide is delayed, and as a result, all of alloy elements contributing to improvement of grain boundary strength are not segregated to grain boundaries, and E _GB is less than 0.50. As a result, hydrogen embrittlement resistance is reduced.

In production No. B-2, since the rolling start temperature of hot rolling is low, unrecrystallized austenite remains, and precipitation of alloy carbide is delayed, and E _GB is less than 0.50. As a result, hydrogen embrittlement resistance is reduced.

In production No. C-2, since the rolling start temperature of hot rolling is high, grains of recrystallized austenite are coarsened, precipitation of alloy carbide is delayed, and E _GB is less than 0.50. As a result, hydrogen embrittlement resistance is reduced.

In production No. D-2, since the rolling end temperature is low, unrecrystallized austenite remains, and precipitation of alloy carbide is delayed, and E _GB is less than 0.50. As a result, hydrogen embrittlement resistance is reduced.

In production No. E-2, since the rolling end temperature is high, grains of recrystallized austenite are coarsened, precipitation of alloy carbide is delayed, and E _GB is less than 0.50. As a result, hydrogen embrittlement resistance is reduced.

In production No. F-2, since the diffusion of the alloy element is delayed due to the low coiling temperature, precipitation of the alloy carbide is suppressed, and E _GB is less than 0.50. As a result, hydrogen embrittlement resistance is reduced.

In production No. G-2, since the coiling temperature is high, an internal oxide layer is formed in the surface layer of the hot-rolled steel sheet, and cracks occur in the surface of the steel sheet in the subsequent treatment. Therefore, analysis of the tissue and evaluation of mechanical properties were not performed.

In production No. H-2, since the average cooling rate to the coiling temperature is low, ferrite and pearlite transformation occurs, precipitation of alloy carbide is suppressed, and E _GB is less than 0.50. As a result, hydrogen embrittlement resistance is reduced.

In the production of No. K-2, since the inter-pass time of the pass having a reduction ratio of 20% or more in finish rolling is long, precipitation of alloy carbide in the hot rolling step is delayed. As a result, E _GB was lower than 0.50, and hydrogen embrittlement resistance was lowered.

In the production of No. L-2, the holding temperature after the hot rolling step is low, the residence time at 400 to 550 ℃ is less than 600 seconds, and precipitation of alloy carbide is not sufficiently generated, so that E _GB is less than 0.50. As a result, the hydrogen embrittlement resistance is deteriorated.

In the production No. M-2, since the holding temperature after the hot rolling step is high and the residence time at 400 to 550℃is not satisfactory for 600 seconds, precipitation of alloy carbide does not sufficiently occur, and E _GB is less than 0.50. As a result, hydrogen embrittlement resistance is reduced.

Production No. A-3 has a low annealing temperature, and thus undergoes ferrite transformation during the holding, and the tensile strength is lower than 1500MPa.

Production No. B-3 has a high annealing temperature, and therefore, the concentration of the alloying element segregated in the grain boundary is reduced. As a result, E _GB was lower than 0.50, and hydrogen embrittlement resistance was lowered.

In the production No. C-3, since the average cooling rate at the self-annealing temperature is low, ferrite transformation occurs during the cooling process, and the tensile strength does not reach 1500MPa.

The production No. F-3 has a high cooling stop temperature, and thus a bainite transformation occurs, and the tensile strength does not reach 1500MPa.

Production No. H-3 has a high tempering temperature, and thus, it produces softening of martensite, and the tensile strength does not reach 1500MPa.

In the production of No. J-3, the tempering time was long, and thus the softening of martensite was excessive, and the tensile strength did not reach 1500MPa.

In the production of No. K-3, since the cooling stop temperature and the intermediate holding temperature in the post-annealing cooling step are low (outside the 2 nd temperature zone), bainite transformation occurs, and the tensile strength does not reach 1500MPa.

In production No. L-3, since the cooling stop temperature and the intermediate holding temperature in the post-annealing cooling step are high (outside the 2 nd temperature zone), ferrite transformation and pearlite transformation occur, and the tensile strength does not reach 1500MPa.

Production No. M-3 has a long holding time at the second cooling stop temperature, and thus a bainite transformation occurs during holding, and the tensile strength does not reach 1500MPa.

Since the cooling rate from the second cooling stop temperature was low, ferrite transformation and bainite transformation were generated during the cooling process, and the tensile strength did not reach 1500MPa in production No. n-3.

In contrast, in all the examples according to the present invention, by properly controlling hot rolling, coiling, annealing, and the like, a steel sheet having high strength and excellent hydrogen embrittlement resistance can be obtained.

Fig. 1 is a graph showing the relationship between E _GB and tensile strength, which affect hydrogen embrittlement resistance of the steel sheets of examples 1 and 2. In FIG. 1, # is an example in which the hydrogen embrittlement resistance does not reach the target, and # is an example in which the hydrogen embrittlement resistance reaches the target. As shown in fig. 1, when E _GB is set to 0.50 or more, excellent hydrogen embrittlement resistance can be obtained even for a high-strength material of 1500MPa or more.

Industrial applicability

According to the present invention, a high-strength steel sheet having excellent hydrogen embrittlement resistance can be provided. When used as a steel sheet for automobiles, the steel sheet contributes to improvement of fuel efficiency by reducing the weight of a vehicle body.

Claims

1.A steel sheet, characterized in that,

The chemical composition of the material comprises the following components in mass percent: c:0.150 to 0.400 percent,

Si：0.01～2.00％、

Mn：0.80～2.00％、

P：0.0001～0.0200％、

S：0.0001～0.0200％、

Al：0.001～1.000％、

N：0.0001～0.0200％、

O：0.0001～0.0200％、

Co：0～0.500％、

Ni：0～1.000％、

Mo：0～1.000％、

Cr：0～2.000％、

Ti：0～0.500％、

B：0～0.0100％、

Nb：0～0.500％、

V：0～0.500％、

Cu：0～0.500％、

W：0～0.100％、

Ta：0～0.100％、

Mg：0～0.050％、

Ca：0～0.050％、

Y：0～0.050％、

Zr：0～0.050％、

La：0～0.050％、

Ce：0～0.050％、

Sn：0～0.050％、

Sb：0～0.050％、

As:0 to 0.050%, and

The remainder: fe and impurities are mixed with each other,

The microstructure comprises in area ratio:

ferrite: less than 5.0 percent,

Martensite and tempered martensite: total more than 90.0%, and

The remainder: 1 or more than 2 kinds of bainite, pearlite and retained austenite,

When the interface between adjacent martensite and tempered martensite having a difference in orientation of 15 DEG or more is used as a prior austenite grain boundary, the bonding strength energy E _GB determined by the concentration of each alloy element in the prior austenite grain boundary satisfies the following formula (1),

The tensile strength is more than 1500 MPa;

E_GB＝1+(3×[Co]+0.7×[Ni]+5.5×[Mo]+0.7×[Cr]+2.9×[Ti]+47×

[B]+4.3×[Nb]+4.5×[V]+5.2×[W]+3.1×[Ta]+4.3×[Zr]-0.25×[Mn]-0.1×[P]-[Cu]-1.1×[Sn]-0.6×[Sb]-0.9×[As])≥0.50 (1)

2. The steel sheet according to claim 1, wherein,

The chemical composition contains 1 or more than 2 elements selected from the following elements:

Co：0.01～0.500％、

Ni：0.01～1.000％、

Mo：0.01～1.000％、

Cr：0.001～2.000％、

Ti：0.001～0.500％、

B：0.0001～0.0100％、

Nb：0.001～0.500％、

V：0.001～0.500％、

Cu：0.001～0.500％、

W：0.001～0.100％、

Ta：0.001～0.100％、

Mg：0.001～0.050％、

Ca：0.001～0.050％、

Y：0.001～0.050％、

Zr：0.001～0.050％、

La：0.001～0.050％、

Ce：0.001～0.050％、

Sn：0.001～0.050％、

Sb:0.001 to 0.050%, and

As：0.001～0.050％。

3. The steel sheet according to claim 1 or 2, wherein the steel sheet has a coating layer comprising zinc, aluminum, magnesium or an alloy thereof on the surface.