EP4600387A1

EP4600387A1 - High-strength steel sheet, member formed using high-strength steel sheet, automobile framework structure component or automobile reinforcing component composed of member, and production methods for high-strength steel sheet and member

Info

Publication number: EP4600387A1
Application number: EP22967883.4A
Authority: EP
Inventors: Yusuke Wada; Hidekazu Minami; Yuki Toji; Yoshiyasu Kawasaki
Original assignee: JFE Steel Corp
Current assignee: JFE Steel Corp
Priority date: 2022-12-08
Filing date: 2022-12-08
Publication date: 2025-08-13
Also published as: JPWO2024122037A1; WO2024122037A1; JP7367893B1; EP4600387A4; MX2025006648A; CN120265805A; KR20250075720A

Abstract

Provided is a high-strength steel sheet excellent in component strength, ductility, stretch flangeability, bendability of sheared end surface portions, and low-temperature toughness. The high-strength steel sheet comprises: a predetermined chemical composition; and a microstructure at a position of 1/4 of a sheet thickness of the high-strength steel sheet in which an area ratio of martensite is 10 % to 80 %, an area ratio of bainite is 2 % to 70 %, an area ratio of ferrite is 80 % or less, an area ratio of retained austenite is 15 % or less, and a proportion of a number of martensite blocks in which metastable carbides are present to that of martensite blocks is 2 % or more, wherein when nanohardness is measured at 225 points or more at the position of 1/4 of the sheet thickness, a standard deviation σ_n of the nanohardness is 0.60 × [H_n]^ave or less.

Description

TECHNICAL FIELD

The present disclosure relates to a high-strength steel sheet, a member produced using the high-strength steel sheet, an automotive framework structural component or automotive reinforcing component comprising the member, and production methods for the high-strength steel sheet and the member.

BACKGROUND

Efforts have been made to increase the strength of steel sheets for automobiles in order to improve crashworthiness while reducing CO₂ emissions by reduction of vehicle weight. Moreover, with the introduction of a series of new regulations, there has been a significant increase in the number of cases where high-strength steel sheets with a tensile strength (TS) of 780 MPa or more are used for main structural components and reinforcing components forming the framework of the automotive cabin (hereafter also referred to as automotive framework structural components) in order to increase the strength of the automotive body.
For example, Patent Literature (PTL) 1 proposes a high-strength steel sheet that has a specified chemical composition and microstructure and has a particle size of iron carbide contained in a low-temperature transformation phase of 500 nm or less.

CITATION LIST

Patent Literature

PTL 1: JP 2008-308717 A

SUMMARY

(Technical Problem)

High-strength steel sheets used in automotive framework structural components are required to have excellent component strength (high impact absorbed energy during a collision), and are required to have a high yield stress (YS) and a high yield ratio (YR = yield stress (YS)/tensile strength (TS)). Since components such as crash boxes have punched end surfaces and bent portions, steel sheets used in such components are required to have good ductility, stretch flangeability, and bendability of sheared end surface portions. Moreover, given the possibility that, when components using high-strength steel sheets with a tensile strength of 780 MPa or more are used in a low-temperature environment, toughness degrades and cracks occur during a collision, steel sheets for automobiles are required to have excellent low-temperature toughness to prevent cracks during a collision when used in a low-temperature environment.
A steel sheet that satisfies all of these properties is still needed in order to increase the ratio of high-strength steel sheets used in automotive components (automotive parts).
It could therefore be helpful to provide a high-strength steel sheet excellent in component strength (part strength), ductility, stretch flangeability, bendability of sheared end surface portions, and low-temperature toughness, together with a production method therefor.
It could also be helpful to provide a member produced using the high-strength steel sheet, together with a production method therefor.
Herein, "high-strength steel sheet" means a steel sheet whose tensile strength (TS) determined by the below-described tensile test is 780 MPa or more.
"Excellent component strength" means that the yield ratio (YR) determined by the below-described tensile test is 55 % or more.
"Excellent ductility" means that the total elongation (El) determined by the below-described tensile test is 10 % or more.
"Excellent stretch flangeability" means that the hole expansion ratio (λ) determined by the below-described hole expanding test is 20 % or more.
"Excellent bendability of sheared end surface portions" means that the ratio (Rs/Rg) of the limit bending radius (Rs/t) determined by the below-described bend test of a sample having a sheared end surface portion to the limit bending radius (Rg/t) determined by the below-described bend test of a sample having a ground end surface portion is 1.50 or less.
"Excellent low-temperature toughness" means that the low-temperature toughness parameter (P) is 3000 or more in the below-described Charpy impact test.

(Solution to Problem)

Upon careful examination, we discovered that the stated object can be achieved by employing the following structure. The present disclosure is based on this discovery.
We thus provide:

(1) A high-strength steel sheet comprising: a chemical composition containing (consisting of), in mass%, C: 0.030 % or more and 0.500 % or less, Si: 0.01 % or more and 2.50 % or less, Mn: 0.10 % or more and 5.00 % or less, P: 0.100 % or less, S: 0.0200 % or less, Al: 1.000 % or less, N: 0.0100 % or less, and O: 0.0100 % or less with a balance consisting of Fe and inevitable impurities; and a steel microstructure at a position of 1/4 of a sheet thickness of the high-strength steel sheet in which an area ratio of martensite is 10 % or more and 80 % or less, an area ratio of bainite is 2 % or more and 70 % or less, an area ratio of ferrite is 80 % or less, an area ratio of retained austenite is 15 % or less, and a proportion of a number of martensite blocks in which metastable carbides are present to a total number of martensite blocks is 2 % or more, wherein when nanohardness is measured at 225 points or more at the position of 1/4 of the sheet thickness, a standard deviation σ_n of the nanohardness is 0.60 × [H_n]^ave or less, where [H_n]^ave is an average value of the nanohardness.
(2) The high-strength steel sheet according to (1), wherein an average number density of the metastable carbides in the martensite blocks in which the metastable carbides are present is 1 × 10⁶/mm² or more.
(3) The high-strength steel sheet according to (1) or (2), wherein the chemical composition further contains, in mass%, at least one element selected from the group consisting of Ti: 0.200 % or less, Nb: 0.200 % or less, V: 0.200 % or less, Ta: 0.10 % or less, W: 0.10 % or less, B: 0.0100 % or less, Cr: 1.00 % or less, Mo: 1.00 % or less, Ni: 1.00 % or less, Co: 0.010 % or less, Cu: 1.00 % or less, Sn: 0.200 % or less, Sb: 0.200 % or less, Ca: 0.0100 % or less, Mg: 0.0100 % or less, REM: 0.0100 % or less, Zr: 0.100 % or less, Te: 0.100 % or less, Hf: 0.10 % or less, and Bi: 0.200 % or less.
(4) The high-strength steel sheet according to any one of (1) to (3), comprising a soft surface layer that is a region of 200 µm or less from a surface of the high-strength steel sheet in a sheet thickness direction and having a Vickers hardness of 85 % or less of a Vickers hardness at the position of 1/4 of the sheet thickness of the high-strength steel sheet, wherein when nanohardness is measured at 300 points or more in a 50 µm × 50 µm region of a sheet plane at each of a position of 1/4 of a depth of the soft surface layer from the surface of the high-strength steel sheet in the sheet thickness direction and a position of 1/2 of the depth of the soft surface layer from the surface of the high-strength steel sheet in the sheet thickness direction, a proportion of a number of measurements in which the nanohardness of the sheet plane at the position of 1/4 of the depth of the soft surface layer from the surface of the high-strength steel sheet in the sheet thickness direction is 7.0 GPa or more is 0.10 or less relative to a total number of measurements, a standard deviation σ of the nanohardness of the sheet plane at the position of 1/4 of the depth of the soft surface layer from the surface of the high-strength steel sheet in the sheet thickness direction is 1.8 GPa or less, and a standard deviation σ of the nanohardness of the sheet plane at the position of 1/2 of the depth of the soft surface layer from the surface of the high-strength steel sheet in the sheet thickness direction is 2.2 GPa or less.
(5) The high-strength steel sheet according to any one of (1) to (4), comprising a metal coated or plated layer on a surface of the high-strength steel sheet on one side or both sides, the metal coated or plated layer containing one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi in a total amount of more than 50 mass%.
(6) The high-strength steel sheet according to any one of (1) to (5), comprising a metal coated or plated layer in an outermost layer of the high-strength steel sheet on one side or both sides, the metal coated or plated layer containing at least one of zinc or aluminum in a total amount of 50 mass% or more.
(7) A member produced using the high-strength steel sheet according to any one of (1) to (6).
(8) An automotive framework structural component or automotive reinforcing component comprising the member according to (7).
(9) A production method for a high-strength steel sheet, the production method comprising: a hot rolling process of subjecting a steel slab having the chemical composition according to (1) or (3) to rough rolling under a condition that an average strain rate is 1 × 10^-4/s or more and 1 × 10^-1/s or less and a total rolling reduction ratio is 50 % or more, thereafter to finish rolling, and thereafter to coiling treatment to obtain a hot-rolled sheet; thereafter a pickling and cold rolling process of subjecting the hot-rolled sheet to pickling and cold rolling to obtain a cold-rolled sheet; thereafter a first heating process of subjecting the cold-rolled sheet to first heating under a condition that a heating temperature is 750 °C or more; thereafter a first cooling process of cooling the cold-rolled sheet under a condition that a first cooling rate in a temperature range of T₂ or more and 750 °C or less is 2.0 °C/s or more; thereafter an in-furnace retention process of retaining the cold-rolled sheet under a condition that a retention temperature T₂ is 350 °C or more and 550 °C or less and a retention time t in seconds satisfies that F defined in the following Formula 1 is 0.20 or more and 0.90 or less, $F = 1 - \exp (- {kt}^{n})$ where t is the retention time in seconds, and k and n are constants determined from an expansion curve of a Formaster test in which a test piece obtained by subjecting the slab to processes up to end of the first cooling process is held at the retention temperature T₂ of 350 °C or more and 550 °C or less; thereafter a second cooling process of cooling the cold-rolled sheet to Ms - 20 °C or less under a condition that a second average cooling rate in a temperature range of Ms - 20 °C or more and Ms or less is 5 °C/s or more; and thereafter a second heating process of treating the cold-rolled sheet under a condition that a temperature X in °C and a holding time Y in seconds satisfy the following Formula 2, $7000 \leq (273 + X) (20 + \log (Y / 3600)) \leq 13000 .$
(10) The production method for a high-strength steel sheet according to (9), wherein in the second heating process, the temperature X in °C satisfies the following Formula 3, $100 \leq X \leq 400 .$
(11) The production method for a high-strength steel sheet according to (9) or (10), wherein the first heating process is performed in an atmosphere with a dew point of -30 °C or more.
(12) The production method for a high-strength steel sheet according to any one of (9) to (11), comprising a process of applying, to one side or both sides of the steel sheet after the cold rolling process and before the annealing process, metal coating or plating containing one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi in an amount of more than 50 mass%.
(13) The production method for a high-strength steel sheet according to any one of (9) to (12), comprising a process of applying, to the steel sheet from the first heating to the second heating process, metal coating or plating containing at least one of zinc or aluminum in a total amount of 50 mass% or more.
(14) A production method for a member, the production method comprising subjecting the high-strength steel sheet according to any one of (1) to (6) to at least one of forming or joining to obtain the member.

(Advantageous Effect)

It is thus possible to provide a high-strength steel sheet excellent in component strength, ductility, stretch flangeability, bendability of sheared end surface portions, and low-temperature toughness, and a member produced using the high-strength steel sheet.
It is also possible to provide production methods for the high-strength steel sheet and the member produced using the high-strength steel sheet.
It is also possible to provide an automotive framework structural component or automotive reinforcing component comprising the member.

BRIEF DESCRIPTION OF THE DRAWINGS

In the accompanying drawings:

FIG. 1 is a diagram illustrating an example of an electron diffraction pattern of martensite in which carbides are present;
FIGS. 2A and 2B are schematic diagrams relating to preparation of a sample for a V bend + orthogonal VDA bend test in examples, where FIG. 2A relates to V bending (primary bending) and FIG. 2B relates to orthogonal VDA bending (secondary bending); and
FIGS. 3A to 3C are schematic diagrams relating to a sample for an axial crush test and to the test in examples, where FIG. 3A is a front view of the test member, FIG. 3B is a perspective view of the test member, and FIG. 3C is a schematic diagram illustrating the axial crush test.

DETAILED DESCRIPTION

The presently disclosed technology will be described by way of embodiments below. The present disclosure is not limited to the following embodiments.

[High-strength steel sheet]

A high-strength steel sheet (hereafter also referred to as "steel sheet" for convenience) according to the present disclosure comprises the below-described chemical composition and steel microstructure.

The chemical composition of the high-strength steel sheet according to the present disclosure (hereafter also referred to as "chemical composition according to the present disclosure" for convenience) will be described. "%" in the chemical composition according to the present disclosure is "mass%" unless otherwise specified.

(C: 0.030 % or more and 0.500 % or less)

C is one of the important basic components of steel. In the present disclosure, in particular, C influences the area ratio of martensite. If the C content is excessively low, the area ratio of martensite decreases, making it difficult to achieve a TS of 780 MPa or more. The C content is therefore 0.030 % or more, preferably 0.040 % or more, and more preferably 0.050 % or more.
If the C content is excessively high, the amount of retained austenite increases excessively, and the hardness of martensite generated from retained austenite during blanking increases significantly. As a result, crack propagation during hole expansion is promoted, the hole expansion ratio decreases, and stretch flangeability decreases. In addition, due to stress-induced transformation of retained austenite, YR decreases and component strength decreases. The C content is therefore 0.500 % or less, preferably 0.400 % or less, and more preferably 0.300 % or less.

(Si: 0.01 % or more and 2.50 % or less)

Si is a component that increases the strength of the steel sheet by suppressing the precipitation of cementite in martensite and by solid solution strengthening. To achieve this effect, the Si content is 0.01 % or more, preferably 0.05 % or more, and more preferably 0.10 % or more.
If the Si content is excessively high, carbide precipitation during bainite transformation is significantly suppressed, the amount of retained austenite increases excessively, and the hardness of martensite generated from retained austenite during blanking increases significantly. As a result, crack propagation during hole expansion is promoted, the hole expansion ratio decreases, and stretch flangeability decreases. In addition, due to stress-induced transformation of retained austenite, YR decreases and component strength decreases. The Si content is therefore 2.50 % or less, preferably 2.00 % or less, and more preferably 1.50 % or less.

(Mn: 0.10 % or more and 5.00 % or less)

Mn is one of the important basic components of steel. In the present disclosure, in particular, Mn influences the area ratio of martensite.
If the Mn content is excessively low, the area ratio of martensite decreases, making it difficult to achieve a TS of 780 MPa or more. The Mn content is therefore 0.10 % or more, preferably 0.90 % or more, and more preferably 1.80 % or more.
If the Mn content is excessively high, austenite stabilizes, the amount of retained austenite increases excessively, and the hardness of martensite generated from retained austenite during blanking increases significantly. As a result, crack propagation during hole expansion is promoted, the hole expansion ratio decreases, and stretch flangeability decreases. In addition, due to stress-induced transformation of retained austenite, YR decreases and component strength decreases. The Mn content is therefore 5.00 % or less, preferably 4.20 % or less, and more preferably 3.60 % or less.

(P: 0.100 % or less)

P is a component that segregates to and embrittles prior austenite grain boundaries and decreases the ultimate deformability of the steel sheet, and thus may cause a decrease in λ and a decrease in bendability. The P content is therefore 0.100 % or less, and preferably 0.070 % or less.
Although no lower limit is placed on the P content, the P content is preferably 0.001 % or more given that P is a solid-solution-strengthening element and can increase the strength of the steel sheet.

(S: 0.0200 % or less)

S is a component that exists as sulfides and decreases the ultimate deformability of the steel sheet, and thus may cause a decrease in λ and a decrease in bendability. The S content is therefore 0.0200 % or less, and preferably 0.0050 % or less.
Although no lower limit is placed on the S content, the S content is preferably 0.0001 % or more due to production technology constraints.

(Al: 1.000 % or less)

Al is an effective component for sufficient deoxidation and reducing inclusions in steel. If the Al content is excessively high, however, a large amount of ferrite is generated and the hole expansion ratio decreases, and stretch flangeability may decrease. The Al content is therefore 1.000 % or less, preferably 0.500 % or less, and more preferably 0.100 % or less.
For stable deoxidation, the Al content is preferably 0.010 % or more, more preferably 0.015 % or more, and further preferably 0.020 % or more.

(N: 0.0100 % or less)

N is a component that exists as nitrides and decreases the ultimate deformability of the steel sheet, and thus may cause a decrease in λ and a decrease in bendability. The N content is therefore 0.0100 % or less, and preferably 0.0050 % or less.
Although no lower limit is placed on the N content, the N content is preferably 0.0001 % or more due to production technology constraints.

(O: 0.0100 % or less)

O is a component that exists as oxides and decreases the ultimate deformability of the steel sheet, and thus may cause a decrease in λ and a decrease in bendability. The O content is therefore 0.0100 % or less, and preferably 0.0050 % or less.
Although no lower limit is placed on the O content, the O content is preferably 0.0001 % or more due to production technology constraints.

(Optional components)

In addition to the foregoing chemical composition, the high-strength steel sheet according to the present disclosure may further contain, in mass%, at least one element selected from the group consisting of

Ti: 0.200 % or less, Nb: 0.200 % or less, V: 0.200 % or less,
Ta: 0.10 % or less, W: 0.10 % or less,
B: 0.0100 % or less,
Cr: 1.00 % or less, Mo: 1.00 % or less, Ni: 1.00 % or less,
Co: 0.010 % or less,
Cu: 1.00 % or less,
Sn: 0.200 % or less,
Sb: 0.200 % or less,
Ca: 0.0100 % or less, Mg: 0.0100 % or less, REM: 0.0100 % or less,
Zr: 0.020 % or less, Te: 0.020 % or less,
Hf: 0.10 % or less, and
Bi: 0.200 % or less.

These elements may be used alone or in combination of two or more.
In the case where any of Ti, Nb, and V is contained, the Ti, Nb, and V contents are each preferably 0.200 % or less and more preferably 0.100 % or less in order to avoid the formation of a large amount of coarse precipitates or inclusions, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Ti, Nb, and V contents, the Ti, Nb, and V contents are each preferably 0.001 % or more because the strength of the steel sheet increases as a result of the formation of fine carbides, nitrides, or carbonitrides during hot rolling or continuous annealing.
In the case where any of Ta and W is contained, the Ta and W contents are each preferably 0.10 % or less and more preferably 0.08 % or less in order to avoid the formation of a large amount of coarse precipitates or inclusions, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Ta and W contents, the Ta and W contents are each preferably 0.01 % or more because the strength of the steel sheet increases as a result of the formation of fine carbides, nitrides, or carbonitrides during hot rolling or continuous annealing.
In the case where B is contained, the B content is preferably 0.0100 % or less and more preferably 0.0003 % or more in order to avoid the occurrence of cracks inside the steel sheet during casting or hot rolling, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the B content, the B content is preferably 0.0003 % or more given that B is an element that segregates to the austenite grain boundaries during annealing and improves quench hardenability.
In the case where any of Cr, Mo, and Ni is contained, the Cr, Mo, and Ni contents are each preferably 1.00 % or less and more preferably 0.80 % or less in order to avoid the increase of coarse precipitates or inclusions, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Cr, Mo, and Ni contents, the Cr, Mo, and Ni contents are each preferably 0.01 % or more given that Cr, Mo, and Ni are elements that improve quench hardenability.
In the case where Co is contained, the Co content is preferably 0.010 % or less and more preferably 0.008 % or less in order to avoid the increase of coarse precipitates or inclusions, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Co content, the Co content is preferably 0.001 % or more given that Co is an element that improves quench hardenability.
In the case where Cu is contained, the Cu content is preferably 1.00 % or less and more preferably 0.80 % or less in order to avoid the increase of coarse precipitates or inclusions, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Cu content, the Cu content is preferably 0.01 % or more given that Cu is an element that improves quench hardenability.
In the case where Sn is contained, the Sn content is preferably 0.200 % or less and more preferably 0.100 % or less in order to avoid the occurrence of cracks inside the steel sheet during casting or hot rolling, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Sn content, the Sn content is preferably 0.001 % or more given that Sn is an element that improves quench hardenability.
In the case where Sb is contained, the Sb content is preferably 0.200 % or less and more preferably 0.100 % or less in order to avoid the increase of coarse precipitates or inclusions, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Sb content, the Sb content is preferably 0.001 % or more given that Sb is an element that controls the thickness of the soft surface layer and enables strength adjustment.
In the case where any of Ca, Mg, and REM is contained, the Ca, Mg, and REM contents are each preferably 0.0100 % or less and more preferably 0.0050 % or less in order to avoid the increase of coarse precipitates or inclusions, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Ca, Mg, and REM contents, the Ca, Mg, and REM contents are each preferably 0.0001 % or more given that Ca, Mg, and REM are elements that make the shape of nitrides and sulfides spheroidal and improve the ultimate deformability of the steel sheet.
In the case where any of Zr and Te is contained, the Zr and Te contents are each preferably 0.100 % or less and more preferably 0.080 % or less in order to avoid the increase of coarse precipitates or inclusions, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Zr and Te contents, the Zr and Te contents are each preferably 0.001 % or more given that Zr and Te are elements that make the shape of nitrides and sulfides spheroidal and improve the ultimate deformability of the steel sheet.
In the case where Hf is contained, the Hf content is preferably 0.10 % or less and more preferably 0.08 % or less in order to avoid the increase of coarse precipitates or inclusions, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Hf content, the Hf content is preferably 0.01 % or more given that Hf is an element that makes the shape of nitrides and sulfides spheroidal and improves the ultimate deformability of the steel sheet.
In the case where Bi is contained, the Bi content is preferably 0.200 % or less and more preferably 0.100 % or less in order to avoid the increase of coarse precipitates or inclusions, which causes the ultimate deformability of the steel sheet to decrease and results in a decrease in λ and a decrease in bendability. Although no lower limit is placed on the Bi content, the Bi content is preferably 0.001 % or more given that Bi is an element that reduces segregation.
The high-strength steel sheet according to one embodiment of the present disclosure has a chemical composition that contains the foregoing essential components and may also contain the foregoing optional components, with the balance consisting of Fe and inevitable impurities. Examples of the inevitable impurities include Zn, Pb, As, Ge, Sr, and Cs. An allowable total content of these inevitable impurities is 0.100 % or less.

The steel microstructure of the high-strength steel sheet according to the present disclosure will be described.

(Area ratio of martensite: 10 % or more and 80 % or less)

By containing martensite, a TS of 780 MPa or more can be easily achieved. The area ratio of martensite is therefore 10 % or more, preferably 15 % or more, and more preferably 20 % or more.
If there is too much martensite, El decreases and ductility degrades. The area ratio of martensite is therefore 80 % or less, preferably 75 % or less, and more preferably 70 % or less.
Herein, martensite includes lower bainite, martensite that has undergone self-tempering during cooling in the below-described annealing, martensite that has been tempered in the below-described second heating process, etc.
The observation position of martensite is a position of 1/4 of the sheet thickness of the steel sheet, as described later.

(Area ratio of bainite: 2 % or more and 70 % or less)

By containing bainite, the difference in hardness between microstructures decreases and λ increases. Moreover, interfacial crack propagation is suppressed to improve low-temperature toughness. The area ratio of bainite is therefore 2 % or more, preferably 3 % or more, and more preferably 4 % or more.
If there is too much bainite, a sufficient amount of martensite cannot be secured and TS decreases. The area ratio of bainite is therefore 70 % or less, preferably 60 % or less, and more preferably 50 % or less.
Herein, bainite is a mixed microstructure of bainitic ferrite with an angular shape, iron-based carbides, and retained austenite that is formed in a temperature range of Ms or more and 700 °C or less.
The observation position of bainite is the position of 1/4 of the sheet thickness of the steel sheet, as described later.

(Area ratio of ferrite: 80 % or less)

By limiting the area ratio of ferrite to 80 % or less, the desired strength can be easily obtained. The effects according to the present disclosure can be achieved even when the area ratio of ferrite is 0 %. If there is too much ferrite, a sufficient amount of martensite cannot be secured and the desired TS cannot be achieved. The area ratio of ferrite is therefore 80 % or less, preferably 75 % or less, and more preferably 70 % or less. To increase El and further improve ductility, the area ratio of ferrite is preferably 10 % or more and more preferably 15 % or more.
Herein, ferrite is soft bcc iron formed at a relatively high temperature, and includes allotriomorphic ferrite and idiomorphic ferrite.
The observation position of ferrite is the position of 1/4 of the sheet thickness of the steel sheet, as described later.
The method of measuring the area ratio of each of martensite, bainite, and ferrite is as follows.
First, a sample is cut out from the steel sheet so that a sheet thickness section (L-section at the position of 1/4 of the sheet thickness) parallel to the rolling direction will be the observation plane. The observation plane of the sample is mirror-polished using diamond paste, then finish-polished using colloidal silica, and then etched with 1 vol% nital to reveal the microstructure.
The observation plane of the sample is then observed using a scanning electron microscope (SEM) at 3000x magnification with an accelerating voltage of 10 kV, and SEM images for three observation fields (40 µm × 30 µm per observation field) are obtained.
The area ratio of each microstructure is calculated from the obtained SEM images using Adobe Photoshop (produced by Adobe Systems Inc.). Specifically, a value obtained by dividing the area of each microstructure by the measurement area is taken as the area ratio of the microstructure. The area ratio of each microstructure is calculated for the three observation fields, and the average value is taken as the area ratio of the microstructure.
In SEM images, ferrite is a gray, flat microstructure region that does not contain carbides with white contrast.
Bainite is a mixed microstructure region composed of gray, angular bainitic ferrite, iron-based carbides with white contrast, and acicular retained austenite.
Martensite is a microstructure having a hierarchical structure with fine irregularities inside.
These are distinguishable from each other.

(Area ratio of retained austenite: 15 % or less)

Reducing retained austenite makes it possible to obtain good component strength and stretch flangeability. The area ratio of retained austenite is therefore 15 % or less, and preferably 10 % or less. No lower limit is placed on the area ratio of retained austenite, and the effect can still be achieved with an area ratio of retained austenite of 0 %.
The method of measuring the area ratio of retained austenite is as follows.
First, the steel sheet is ground so that the position of 1/4 of the sheet thickness (i.e. position corresponding to 1/4 of the sheet thickness from the steel sheet surface in the depth direction) will be the measurement plane, and then further polished by 0.1 mm by chemical polishing to obtain a sample.
For the measurement plane of the sample, the integrated reflection intensity of each of (200) plane, (220) plane, and (311) plane of fcc iron (austenite) and (200) plane, (211) plane, and (220) plane of bcc iron is measured using an X-ray diffractometer with a Co Kα radiation source.
The intensity ratio of the integrated reflection intensity of each plane of fcc iron to the integrated reflection intensity of each plane of bcc iron is calculated. The average of the nine intensity ratios is taken as the volume ratio of retained austenite. Assuming that the volume ratio of retained austenite is three-dimensionally uniform, the volume ratio of retained austenite is taken as the area ratio of retained austenite at the position of 1/4 of the sheet thickness of the steel sheet.

(Residual microstructure)

The steel microstructure according to the present disclosure may contain microstructure (residual microstructure) other than the foregoing martensite, bainite, ferrite, and retained austenite.
The residual microstructure includes microstructures that are other than martensite, bainite, ferrite, and retained austenite and are known as microstructures of steel sheets, such as pearlite and alloy carbonitrides precipitated in ferrite.
The residual microstructure does not include iron-based carbides present in bainite, metastable carbides precipitated in martensite, and iron-based carbides such as cementite precipitated in martensite.
The area ratio of the residual microstructure is preferably 3 % or less because the effects according to the present disclosure are not undermined.

(Proportion of the number of martensite blocks in which metastable carbides are present to the number of martensite blocks: 2 % or more)

Metastable carbides precipitated in martensite blocks improve low-temperature toughness while maintaining excellent component strength, ductility, sheared end surface bendability, and stretch flangeability. This is considered to be because metastable carbides precipitated in martensite blocks suppress the initiation and propagation of cracks at low temperatures.
To achieve this effect, the proportion of the number of martensite blocks in which metastable carbides are present to the number of martensite blocks (hereafter also referred to as "proportion p") is 2 % or more, preferably 5 % or more, more preferably 10 % or more, further preferably 20 % or more, and particularly preferably 30 % or more. No upper limit is placed on the proportion p, and the upper limit may be 100 %.
Herein, metastable carbides are carbides that are metastable and precipitate during the process of tempering martensite. Metastable carbides are, for example, Fe carbides (iron-based carbides) other than cementite, and include at least one type of carbide selected from the group consisting of epsilon (ε) carbide, eta (η) carbide, and chi (χ) carbide.
The method of measuring the proportion of the number of martensite blocks in which metastable carbides are present to the number of martensite blocks (proportion p) is as follows.
First, the steel sheet is ground so that the position of 1/4 of the sheet thickness (i.e. position corresponding to 1/4 of the sheet thickness from the steel sheet surface in the depth direction) will be the observation plane, and then electropolished to prepare a sample. The observation plane of the prepared sample is observed using a transmission electron microscope (TEM) with an accelerating voltage of 200 kV.
When an electron beam is incident on a martensite block from [100] orientation, an electron diffraction pattern of matrix phase martensite is obtained. Adjacent martensite blocks have different crystal orientations across block boundaries and thus differ in contrast in bright field images, and therefore can be distinguished from each other. Martensite can be distinguished from ferrite and bainite as high-density dislocations are observed in martensite whereas the dislocation density is relatively low in ferrite and bainite.
FIG. 1 illustrates an example of an electron diffraction pattern of martensite in which carbides are present. In the case where carbides are present in a single martensite block observed, an electron diffraction pattern of carbides is obtained in addition to an electron diffraction pattern of matrix phase martensite (α), as illustrated in FIG. 1.
In FIG. 1, the black circles indicate electron diffraction spots of matrix phase martensite and the white circles indicate electron diffraction spots of carbides when an electron beam is incident from [100] orientation.
From the distance D1 between the electron diffraction spots of matrix phase martensite and the distance D2 between the electron diffraction spots of carbides in FIG. 1, the ratio dc/dm of the interplanar spacing dc of carbides to the interplanar spacing dm of matrix phase martensite is calculated using the following Formula 3. $dc / dm = D 1 / D 2$
If the ratio dc/dm of the interplanar spacing dc of carbides to the interplanar spacing dm of matrix phase martensite is 1.020 or more and 1.150 or less in the observed martensite block, the martensite block is defined as a martensite block in which metastable carbides are present. In other words, the proportion of the number of martensite blocks in which metastable carbides are present to the number of martensite blocks is the proportion of the number of martensite blocks having a ratio dc/dm of 1.020 or more and 1.150 or less to the number of martensite blocks.
The distance D1 between the electron diffraction spots of matrix phase martensite is a constant value, whereas the distance D2 between the electron diffraction spots of carbides varies depending on the type of carbide.
For example, if the carbide is cementite, the distance D2 is equal to the distance D1. Hence, the ratio dc/dm is 1 (D1/D2 = dc/dm = 1).
Meanwhile, the distance D2 of metastable carbides (such as ε carbide) is shorter than that of cementite, so that the ratio dc/dm (= D1/D2) is greater than 1. Accordingly, the lower limit of the ratio dc/dm when metastable carbides are present is set to 1.020.
The upper limit of the ratio dc/dm is set to 1.150 based on the distance D2 of ε carbide.
Metastable carbides may be present inside martensite blocks or at boundary portions such as block boundaries, but are preferably present inside martensite blocks.
50 martensite blocks are observed, and the number of martensite blocks in which metastable carbides are present is divided by the number of martensite blocks observed. The resultant value (the number of martensite blocks in which metastable carbides are present/50) is multiplied by 100 to yield the proportion of the number of martensite blocks in which metastable carbides are present to the number of martensite blocks (proportion p (%)).

(Average number density of metastable carbides in martensite blocks in which metastable carbides are present: 1 × 10⁶/mm² or more)

The number density of metastable carbides in martensite blocks is preferably high from the viewpoint of achieving better low-temperature toughness. It is considered that, if the number density of metastable carbides is high, the resistance to crack propagation in martensite at low temperatures increases.
Specifically, the average number density of metastable carbides in the martensite blocks in which metastable carbides are present (hereafter also referred to as "number density n") is preferably 1 × 10⁶/mm² or more, more preferably 10 × 10⁶/mm² or more, and further preferably 100 × 10⁶/mm² or more.
No upper limit is placed on the number density n. For example, the number density n may be 10000000 × 10⁶/mm² or less, is preferably 1000000 × 10⁶/mm² or less, more preferably 100000 × 10⁶/mm² or less, and further preferably 10000 × 10⁶/mm² or less.
The method of measuring the average number density of metastable carbides in the martensite blocks in which metastable carbides are present (number density n) is as follows.
When measuring the proportion p using a TEM as described above, a selected-area electron diffraction pattern is obtained for a single martensite block in which metastable carbides are present, and a dark field image is obtained using the electron diffraction spots of metastable carbides. In dark field images, metastable carbides exhibit white contrast.
An image of a 300 nm × 300 nm region is taken inside the single martensite block, and the number of metastable carbides is counted. Adjacent martensite blocks may exist in the 300 nm × 300 nm region across block boundaries.
The area of the martensite block in which metastable carbides are present is defined as the area of the single martensite block from which the selected-area electron diffraction pattern is obtained. Adjacent martensite blocks have different crystal orientations across block boundaries and thus differ in contrast in bright field images, and therefore can be distinguished from each other.
The foregoing measurement is carried out for three observation fields, and the number of metastable carbides divided by the area of the martensite block in which metastable carbides are present (= the number of metastable carbides/the area of the martensite block in which metastable carbides are present) is calculated for each of the three observation fields. The average of these values is taken as the average number density of metastable carbides in the martensite blocks in which metastable carbides are present (number density n).

(Average equivalent circular diameter of metastable carbides: 20 nm or less)

When the average equivalent circular diameter of metastable carbides in martensite blocks is smaller, cracks are less likely to occur in martensite at low temperatures, and accordingly low-temperature toughness is better. The average equivalent circular diameter of metastable carbides in martensite blocks is therefore preferably 20 nm or less and more preferably 5 nm or less.
The method of measuring the average equivalent circular diameter of metastable carbides in martensite blocks is as follows.
When measuring the proportion p using a TEM as described above, a selected-area electron diffraction pattern is obtained for a single martensite block in which metastable carbides are present, and a dark field image is obtained using the electron diffraction spots of metastable carbides. In dark field images, metastable carbides exhibit white contrast.
A dark field image of a 300 nm × 300 nm region is taken inside the single martensite block, and image processing is performed to obtain a binary image so that metastable carbides can be distinguished. The binary image is subjected to particle analysis to determine the equivalent circular diameter for each of the metastable carbide particles. If metastable carbides overlap each other in the dark field image, the binary image is divided using the watershed method.
The equivalent circular diameter is obtained for all metastable carbides present in the 300 nm × 300 nm region (for three observation fields). The average value of the equivalent circular diameters for the three observation fields is calculated and taken as the average equivalent circular diameter of metastable carbides in martensite blocks.

By lowering the standard deviation σ_n of the nanohardness of the steel sheet, the sheared end surface bendability is improved while maintaining excellent component strength, ductility, stretch flangeability, and low-temperature toughness. This is considered to be because, as a result of plastic deformation resistance in local regions being made uniform within the microstructure, non-uniformity of plastic deformation in the sheared portion is suppressed and bending deformability in the sheared portion is improved.
With hardness testing methods other than the nanoindentation method, such as the Vickers hardness test, it is impossible to obtain the plastic deformation resistance in local regions of the microstructure at the submicron level. The use of the nanoindentation method provides a solution in the present disclosure.
To achieve the foregoing effect, the standard deviation σ_n of nanohardness is 0.60 × [H_n]^ave or less and preferably 0.50 × [H_n]^ave or less, where [H_n]^ave is the average value of nanohardness (average nanohardness). No lower limit is placed on the standard deviation σ_n of nanohardness, and the lower limit may be 0.
The average nanohardness [Hn]^ave is preferably 3.0 GPa or more. The average nanohardness [Hn]^ave is preferably 9.0 GPa or less. The average nanohardness [Hn]^ave is more preferably 3.5 GPa or more. The average nanohardness [Hn]^ave is more preferably 8.5 GPa or less.
The method of measuring the standard deviation σ_n of nanohardness will be described. A nanoindentation device equipped with a Berkovich indenter is used to determine nanohardness.
A sample is cut out so that a sheet thickness section (L-section) parallel to the rolling direction of the steel sheet will be the observation plane. The observation plane is then mirror-polished using diamond paste, and then finish-polished using colloidal silica. Using the nanoindentation device, nanohardness is measured at 225 points or more on the sample under load control under the conditions of a loading and unloading speed of 50 µN/s, a maximum load of 500 µN, and a data collection pitch of 5 msec. In the measurement, the measurement position is the position of 1/4 of the sheet thickness from the surface of the high-strength steel sheet, and the distance between indentations is 2 µm or more.
A histogram is created from the nanohardness measurement results at 225 points or more, and the standard deviation is calculated and taken as the standard deviation σ_n of nanohardness. The average value of the nanohardness measurement results at 225 points or more is defined as [H_n]^ave.

It is preferable that a soft surface layer is formed in a surface layer of a base steel sheet in the high-strength steel sheet. The soft surface layer contributes to suppressing bending crack propagation during press forming and automotive body collision, with it being possible to improve bending fracture resistance.
Herein, the term "base steel sheet" refers to a high-strength steel sheet that is the base of various coatings or platings in the case of steel sheets subjected to coating or plating treatment, such as hot-dip galvanized steel sheets, galvannealed steel sheets, electrogalvanized steel sheets, and other metal coated or plated steel sheets, and is a high-strength steel sheet in the case of steel sheets not subjected to coating or plating treatment.
The term "surface layer" refers to a region with a thickness of 200 µm from the surface of the base steel sheet to a depth of 200 µm in the sheet thickness direction.
The term "soft layer" refers to a region having a Vickers hardness of 85 % or less of the Vickers hardness of a section (plane parallel to the steel sheet surface) at the position of 1/4 of the sheet thickness of the base steel sheet. The soft layer includes a decarburized layer in the surface layer of the base steel sheet.
The term "soft surface layer" refers to the soft layer included in the surface layer. The whole or part of the surface layer may be the soft layer. The soft surface layer can be a region with a thickness of 200 µm or less from the surface of the base steel sheet in the sheet thickness direction.
For example, suppose a region having a Vickers hardness of 85 % or less of the section (plane parallel to the steel sheet surface) at the position of 1/4 of the sheet thickness of the base steel sheet is formed with a certain depth from the surface of the base steel sheet in the sheet thickness direction. If the certain depth is within 200 µm in the sheet thickness direction, the region with the thickness from the surface to the certain depth in the sheet thickness direction is the soft surface layer. If the certain depth is more than 200 µm in the sheet thickness direction, the region with a thickness of 200 µm from the surface of the base steel sheet to a depth of 200 µm in the sheet thickness direction is the soft surface layer.
In the case where the high-strength steel sheet includes the soft surface layer, no lower limit is placed on the thickness of the soft surface layer, and the thickness of the soft surface layer is preferably 8 µm or more and more preferably more than 17 µm.
Vickers hardness is measured based on JIS Z 2244-1 (2020) with a load of 10 gf.
To achieve excellent bendability during press forming and excellent bending fracture resistance during a collision, when nanohardness is measured at 300 points or more in a 50 µm × 50 µm region of the sheet plane at a position of 1/4 of the depth of the soft surface layer from the surface of the base steel sheet in the sheet thickness direction (position of 1/4 of the thickness of the soft surface layer from the surface of the base steel sheet in the depth direction), the proportion of nanohardness of 7.0 GPa or more is preferably 0.10 or less. If the proportion of nanohardness of 7.0 GPa or more is 0.10 or less, the proportion of hard microstructures (martensite, etc.), inclusions, etc. is low, so that the formation and connection of voids and the propagation of cracks in hard microstructures (martensite, etc.), inclusions, etc. during press forming and collision (crash) can be further suppressed. Thus, excellent R/t and SF_max are obtained.
In the present disclosure, to achieve excellent bendability during press forming and excellent bending fracture resistance during a collision, it is preferable that the standard deviation σ of the nanohardness of the sheet plane at the position of 1/4 of the depth of the soft surface layer from the surface of the base steel sheet in the sheet thickness direction is 1.8 GPa or less and the standard deviation σ of the nanohardness of the sheet plane at the position of 1/2 of the depth of the soft surface layer from the surface of the base steel sheet in the sheet thickness direction is 2.2 GPa or less. If the standard deviation σ of the nanohardness of the sheet plane at the position of 1/4 of the depth of the soft surface layer from the surface of the base steel sheet in the sheet thickness direction is 1.8 GPa or less and the standard deviation σ of the nanohardness of the sheet plane at the position of 1/2 of the depth of the soft surface layer from the surface of the base steel sheet in the sheet thickness direction is 2.2 GPa or less, the difference in microscopic microstructure hardness is small, so that the formation and connection of voids and the propagation of cracks during press forming and collision can be further suppressed. Thus, excellent R/t and SF_max are obtained.
The standard deviation σ of the nanohardness of the sheet plane at the position of 1/4 of the depth of the soft surface layer from the surface of the base steel sheet in the sheet thickness direction is more preferably 1.7 GPa or less. The standard deviation σ of the nanohardness of the sheet plane at the position of 1/2 of the depth of the soft surface layer from the surface of the base steel sheet in the sheet thickness direction is more preferably 2.1 GPa or less.
The nanohardness of the sheet plane at each of the positions of 1/4 and 1/2 of the depth in the sheet thickness direction is hardness measured by the following method.
First, in the case where a coated or plated layer is formed, after peeling off the coated or plated layer, mechanical polishing is performed to the position of 1/4 of the depth of the soft surface layer from the surface of the base steel sheet in the sheet thickness direction. Buffing with diamond and alumina is then performed, and further colloidal silica polishing is performed. The nanohardness is measured using a Berkovich diamond indenter under the following conditions: load: 500 µN, measurement region: 50 µm × 50 µm, and indentation spacing: 2 µm.
Moreover, mechanical polishing is performed to the position of 1/2 of the depth of the soft surface layer in the sheet thickness direction. Buffing with diamond and alumina is then performed, and further colloidal silica polishing is performed. The nanohardness is measured using a Berkovich diamond indenter under the following conditions: load: 500 µN, measurement region: 50 µm × 50 µm, and indentation spacing: 2 µm.
The thickness of the soft surface layer can be measured by the following method. After smoothing a sheet thickness section (L-section) parallel to the rolling direction of the base steel sheet by wet polishing, measurement is performed using a Vickers hardness meter with a load of 10 gf from a position of 1 µm in the sheet thickness direction from the surface of the base steel sheet to a position of 100 µm in the sheet thickness direction at 1 µm intervals. After this, measurement is performed at 20 µm intervals up to the sheet thickness center. A region where the hardness is 85 % or less of the hardness at the position of 1/4 of the sheet thickness is defined as the soft layer (soft surface layer), and the thickness of the region in the sheet thickness direction is taken as the thickness of the soft layer.

The high-strength steel sheet according to one embodiment of the present disclosure preferably includes a first coated or plated layer, which is a metal coated or plated layer, on the surface of the base steel sheet on one side or both sides. The first coated or plated layer is formed directly on the surface of the base steel sheet, and is a metal coated or plated layer containing one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi in a total amount of more than 50 mass%, excluding galvanized layers such as hot-dip galvanized layers, galvannealed layers, and electrogalvanized layers, and hot-dip aluminum-coated layers. The first coated or plated layer is preferably a metal electroplated layer. In the following, a metal electroplated layer will be described as an example.
As a result of the metal electroplated layer being formed on the surface of the steel sheet, the outermost metal electroplated layer contributes to suppressing bending cracks during press forming and automotive body collision, with it being possible to further improve bending fracture resistance.
The metal species of the metal electroplated layer may be any of Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi, but is preferably Fe. In the following, a Fe-based electroplated layer will be described as an example.
The coating weight of the Fe-based electroplated layer is more than 0 g/m², and preferably 2.0 g/m² or more. Although no upper limit is placed on the coating weight of the Fe-based electroplated layer per side, the coating weight of the Fe-based electroplated layer per side is preferably 60 g/m² or less from the viewpoint of cost. The coating weight of the Fe-based electroplated layer is preferably 50 g/m² or less, more preferably 40 g/m² or less, and further preferably 30 g/m² or less.
The coating weight of the Fe-based electroplated layer is measured as follows. A sample of 10 mm × 15 mm size is collected from the Fe-based electroplated steel sheet and embedded in resin to prepare a cross-section resin-embedded sample. Any three parts of the cross section are observed using a scanning electron microscope (SEM) with an accelerating voltage of 15 kV and 2000x to 10000x magnification depending on the thickness of the Fe-based plated layer. The average value of the thicknesses in the three observation fields is multiplied by the specific gravity of iron to calculate the coating weight of the Fe-based plated layer per side.
Examples of the Fe-based electroplated layer include not only pure Fe but also alloy plated layers such as Fe-B alloy, Fe-C alloy, Fe-P alloy, Fe-N alloy, Fe-O alloy, Fe-Ni alloy, Fe-Mn alloy, Fe-Mo alloy, and Fe-W alloy. The chemical composition of the Fe-based electroplated layer is not limited, but preferably contains one or more elements selected from the group consisting of B, C, P, N, O, Ni, Mn, Mo, Zn, W, Pb, Sn, Cr, V, and Co in a total amount of 10 mass% or less, with the balance consisting of Fe and inevitable impurities. If the total content of elements other than Fe is 10 mass% or less, a decrease in electrolysis efficiency can be prevented and the Fe-based electroplated layer can be formed at low costs. In the case of Fe-C alloy, the C content is preferably 0.08 mass% or less.

The high-strength steel sheet according to one embodiment of the present disclosure may include a second coated or plated layer, which is a metal coated or plated layer, as the outermost layer on one side or both sides of the high-strength steel sheet. The second coated or plated layer contains at least one of zinc or aluminum in a total amount of 50 mass% or more, and may be a hot-dip galvanized layer, a galvannealed layer, an electrogalvanized layer, a hot-dip aluminum-coated layer, or the like.
The second coated or plated layer may be formed directly on the surface of the base steel sheet on one side or both sides, or formed on the first coated or plated layer.
Herein, the hot-dip galvanized layer, galvannealed layer, and electrogalvanized layer refer to coated or plated layers containing Zn (zinc) as the main component (specifically, the Zn content is 50.0 mass% or more).
The coated or plated layer of the aluminum coated or plated steel sheet refers to a coated or plated layer containing Al (aluminum) as the main component (specifically, the Al content is 50.0 mass% or more).
For example, the hot-dip galvanized layer preferably contains Zn, 20.0 mass% or less of Fe, and 0.001 mass% or more and 1.0 mass% or less of Al. The hot-dip galvanized layer may optionally contain one or more elements selected from the group consisting of Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, Ca, Cu, Li, Ti, Be, Bi, and REM in a total amount of more than 0.0 mass% and 3.5 mass% or less. The Fe content in the hot-dip galvanized layer is more preferably less than 7.0 mass%. The balance other than the foregoing elements consists of inevitable impurities.
The galvannealed layer preferably contains, for example, 20 mass% or less of Fe and 0.001 mass% or more and 1.0 mass% or less of Al. The galvannealed layer may optionally contain one or more elements selected from the group consisting of Pb, Sb, Si, Sn, Mg, Mn, Ni, Cr, Co, Ca, Cu, Li, Ti, Be, Bi, and REM in a total amount of more than 0 mass% and 3.5 mass% or less. The Fe content in the galvannealed layer is more preferably 7.0 mass% or more and further preferably 8.0 mass% or more. The Fe content in the galvannealed layer is more preferably 15.0 mass% or less and further preferably 12.0 mass% or less. The balance other than the foregoing elements consists of inevitable impurities.
The coating weight of the galvanized layer per side is not limited, but is preferably 20 g/m² or more and 80 g/m² or less.
The coating weight of the galvanized layer is measured as follows. A treatment solution is prepared by adding 0.6 g of a corrosion inhibitor for Fe (IBIT^® 700BK (IBIT is a registered trademark in Japan, other countries, or both) produced by Asahi Chemical Co., Ltd.) to 1 L of a 10 mass% hydrochloric acid aqueous solution. A sample of the steel sheet with the galvanized layer is then immersed in the treatment solution to dissolve the galvanized layer. The mass loss of the sample before and after dissolution is measured, and the value is divided by the surface area of the steel sheet as the base (the surface area of the coated portion) to calculate the coating weight (g/m²).

The sheet thickness of the high-strength steel sheet is not limited, and may be 0.3 mm or more and 3.0 mm or less.

[Production method for high-strength steel sheet]

Next, a production method for a high-strength steel sheet according to the present disclosure (hereafter also referred to as "production method according to the present disclosure" for convenience) will be described. The production method according to the present disclosure is a method of producing the above-described high-strength steel sheet according to the present disclosure. Here, the temperatures in the production method are based on the surface temperature of the steel slab or steel sheet unless otherwise specified.

In the production method according to the present disclosure, first, a steel slab having the above-described chemical composition according to the present disclosure is subjected to hot rolling, pickling, and cold rolling to obtain a cold-rolled sheet.

(Production of steel slab)

For example, a steel slab may be produced by melting steel material to obtain molten steel having the chemical composition according to the present disclosure and solidifying the obtained molten steel. The steelmaking method is not limited, and any known steelmaking method such as converter or electric furnace may be used. The method of producing the steel slab from the molten steel is not limited, and any known method such as continuous casting, ingot casting, or thin slab casting may be used. Continuous casting is preferable from the viewpoint of preventing macrosegregation.

(Hot rolling process)

The produced steel slab is, for example, cooled to room temperature, then reheated and hot rolled (rough rolling and finish rolling), and then coiled. A hot-rolled sheet is thus obtained. The produced steel slab may be, without being cooled to room temperature, charged into a heating furnace as a warm slab, or may be subjected to a short period of heat retention and then immediately subjected to rough rolling.

(Rough rolling)

The steel slab is subjected to rough rolling under the following conditions to obtain a rough-rolled sheet.
The steel slab heating temperature (slab heating temperature) is preferably 1100 °C or more from the viewpoint of melting carbides and reducing the rolling load. To prevent an increase of scale loss, the slab heating temperature is preferably 1300 °C or less. The steel slab heated to the slab heating temperature is subjected to rough rolling.
By optimizing the average strain rate and the total rolling reduction ratio during rough rolling, the standard deviation of nanohardness can be reduced to 0.60 × [H_n]^ave or less. This is considered to be because, in a process of plastic deformation and dynamic recrystallization of austenite grains during rough rolling, solute atoms such as Si and Mn rapidly diffuse through dislocations and grain boundaries of recrystallized grains and thus are appropriately distributed, resulting in uniform plastic deformation resistance in local regions.
Herein, the average strain rate during rough rolling is defined as a value (ε/t_R) obtained by dividing the total rolling ratio ε(-) from the first mill to the final mill of rough rolling by the time t_R (s) required from the start of rolling at the first mill to the completion of rolling at the final mill in rough rolling.
If the average strain rate during rough rolling is more than 1 × 10^-1/s or the total rolling reduction ratio during rough rolling is less than 50 %, the diffusion of solute atoms such as Si and Mn during plastic deformation and dynamic recrystallization of austenite grains is insufficient and the standard deviation of nanohardness exceeds 0.60 × [H_n]^ave.
If the average strain rate during rough rolling is less than 1 × 10^-4/s, the recovery of dislocations in austenite grains is promoted and the recrystallization driving force decreases to thus suppress dynamic recrystallization, so that the diffusion of solute atoms such as Si and Mn is insufficient and the standard deviation of nanohardness exceeds 0.60 × [H_n]^ave.
Accordingly, rough rolling is performed under the conditions that the average strain rate is 1 × 10^-4/s or more and 1 × 10^-1/s or less and the total rolling reduction ratio is 50 % or more. The average strain rate during rough rolling is preferably 1 × 10^-3/s or more and 1 × 10^-2/s or less. The total rolling reduction ratio during rough rolling is preferably 60 % or more.
The rough rolling end temperature is preferably 950 °C or more from the viewpoint of completing the recrystallization of austenite grains. The rough rolling end temperature may be, for example, 1250 °C or less.
If the slab heating temperature is low, it is preferable to heat the rough-rolled sheet using a bar heater or the like before finish rolling from the viewpoint of preventing troubles during hot rolling.

(Finish rolling)

The temperature when performing finish rolling (finish rolling temperature) is preferably the Ar₃ transformation point or more. This reduces the rolling load. Moreover, the rolling reduction ratio in the non-recrystallized state of austenite decreases and the development of abnormal microstructures extending in the rolling direction is suppressed, contributing to excellent workability.
Finish rolling may be performed continuously with rough-rolled sheets joined together. The rough-rolled sheets may be temporarily coiled before finish rolling.
At least part of finish rolling may be conducted as lubrication rolling to reduce the rolling load. Lubrication rolling is preferable from the viewpoint of making the shape and material properties of the steel sheet uniform. In lubrication rolling, the frictional coefficient is preferably in the range of 0.10 or more and 0.25 or less.

(Coiling)

After finish rolling, coiling is performed to obtain a hot-rolled sheet. The coiling temperature after hot rolling is preferably 300 °C or more and 700 °C or less from the viewpoint of improving the sheet passing properties during cold rolling and annealing described later.

(Pickling and cold rolling process)

The hot-rolled sheet obtained by hot rolling is pickled. Pickling removes oxides from the surface of the hot-rolled sheet, as a result of which excellent chemical convertibility, coated or plated layer quality, etc. can be obtained in the high-strength steel sheet as a finished product. Pickling may be performed once or a plurality of times.
The hot-rolled sheet after pickling is optionally subjected to softening heat treatment, and then subjected to cold rolling. A cold-rolled sheet is thus obtained. The cold rolling conditions are not limited, but the total rolling reduction ratio of cold rolling is preferably 20 % or more and 75 % or less. The number of rolling passes and the rolling reduction ratio of each pass are not limited.

(First coating or plating process (optional))

In one embodiment of the present disclosure, the production method may include a first coating or plating process in which a first coated or plated layer, which is a metal coated or plated layer, is formed on one side or both sides of the steel sheet after the hot rolling process (after the cold rolling process if cold rolling is performed) and before the annealing process. The first coating or plating process is preferably a metal electroplating process.
For example, the surface of the cold-rolled sheet obtained as described above may be subjected to metal coating or plating treatment (first coating or plating treatment) such as metal electroplating treatment to obtain a pre-annealing metal coated or plated steel sheet having a pre-annealing metal coated or plated layer formed on at least one side. The metal coated or plated layer referred to here can be the foregoing first coated or plated layer. The pre-annealing metal coated or plated steel sheet is preferably a pre-annealing metal electroplated steel sheet including a pre-annealing metal electroplated layer.
The metal electroplating treatment method is not limited, but it is preferable to perform metal electroplating treatment because the metal coated or plated layer formed on the base steel sheet is preferably a metal electroplated layer as mentioned above. For example, a sulfuric acid bath, a hydrochloric acid bath, or a mixture thereof may be used as a Fe-based electroplating bath. The coating weight of the pre-annealing metal electroplated layer can be adjusted by the current passage time, etc. The "pre-annealing metal electroplated steel sheet" means that the metal electroplated layer has not been subjected to an annealing process, and does not exclude the case where the hot-rolled sheet, hot-rolled and pickled sheet, or cold-rolled sheet before the metal electroplating treatment has been annealed in advance.
The metal species of the electroplated layer may be any of Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi, but is preferably Fe. Hence, a method of Fe-based electroplating will be described below.
The Fe ion content in the Fe-based electroplating bath before current passage start is preferably 0.5 mol/L or more as Fe^{2 +} content. If the Fe ion content in the Fe-based electroplating bath is 0.5 mol/L or more as Fe^{2 +} content, sufficient Fe coating weight can be obtained. The Fe ion content in the Fe-based electroplating bath before current passage start is preferably 2.0 mol/L or less, in order to obtain sufficient Fe coating weight.
The Fe-based electroplating bath may contain at least one element selected from the group consisting of B, C, P, N, O, Ni, Mn, Mo, Zn, W, Pb, Sn, Cr, V, and Co, in addition to Fe ion. The total content of these elements in the Fe-based electroplating bath is preferably set so that the total content of these elements in the pre-annealing Fe-based electroplated layer will be 10 mass% or less. A metal element may be contained as a metal ion, and a non-metal element may be contained as part of boric acid, phosphoric acid, nitric acid, organic acid, and the like. An iron sulfate plating solution may contain a conductivity aid such as sodium sulfate or potassium sulfate, a chelator, and a pH buffering agent.
The other conditions for the Fe-based electroplating bath are not limited. The temperature of the Fe-based electroplating solution is preferably 30 °C or more and preferably 85 °C or less, from the viewpoint of holding at constant temperature. The pH of the Fe-based electroplating bath is not limited. The pH of the Fe-based electroplating bath is preferably 1.0 or more from the viewpoint of preventing a decrease in current efficiency due to hydrogen generation. The pH of the Fe-based electroplating bath is preferably 3.0 or less from the viewpoint of the electric conductivity of the Fe-based electroplating bath. The current density is preferably 10 A/dm² or more from the viewpoint of productivity. The current density is preferably 150 A/dm² or less from the viewpoint of easily controlling the coating weight of the Fe-based electroplated layer. The sheet passing speed is preferably 5 mpm or more from the viewpoint of productivity. The sheet passing speed is preferably 150 mpm or less from the viewpoint of stably controlling the coating weight.
Examples of treatment that may be performed prior to the Fe-based electroplating treatment include degreasing treatment and water washing for cleaning the surface of the cold-rolled sheet, and pickling treatment and water washing for activating the surface of the cold-rolled sheet. Such preliminary treatment is followed by the Fe-based electroplating treatment. The method of degreasing treatment and water washing is not limited, and may be a typical method.
In the pickling treatment, any of various types of acids such as sulfuric acid, hydrochloric acid, nitric acid, and mixtures thereof may be used. Of these, sulfuric acid, hydrochloric acid, or a mixture thereof is preferable. The acid concentration is not limited, but is preferably 1 mass% or more and 20 mass% or less from the viewpoint of, for example, the oxide coating removal ability and the prevention of surface roughening (surface defects) caused by excessive pickling.
The pickling treatment solution may contain a defoamer, a pickling accelerator, a pickling inhibitor, and the like.

(First heating process)

The obtained cold-rolled sheet is then subjected to first heating at 750 °C or more. The cold-rolled sheet may or may not have been subjected to electroplating treatment.
If the first heating temperature is excessively low, reverse transformation to austenite does not proceed sufficiently and the area ratio of martensite decreases, so that the desired TS cannot be obtained. The first heating temperature is therefore 750 °C or more, and preferably 770 °C or more.
Although no upper limit is placed on the heating temperature, the heating temperature is preferably 950 °C or less from the viewpoint of operability, etc.
The time (heating time) for heating the cold-rolled sheet at the first heating temperature is not limited. If the heating time is excessively short, however, reverse transformation to austenite may not proceed sufficiently. The heating time is therefore preferably 30 s or more and more preferably 60 s or more.
No upper limit is placed on the heating time. For example, the heating time may be 6000 s or less, and is preferably 3000 s or less. Here, "s" denotes seconds.
The dew point of the annealing atmosphere in the first heating is preferably -30 °C or more. As a result of the dew point of the annealing atmosphere in the annealing process being -30 °C or more, the decarburization reaction is promoted and the soft surface layer is formed deeper. Consequently, when nanohardness is measured at 300 points or more in a 50 µm × 50 µm region of the sheet plane at the position of 1/4 of the depth of the soft surface layer from the surface of the base steel sheet in the sheet thickness direction, the proportion of nanohardness of 7.0 GPa or more is 0.10 or less. The dew point of the annealing atmosphere in the annealing process is more preferably -15 °C or more and further preferably -5 °C or more. Although no upper limit is placed on the dew point of the annealing atmosphere in the annealing process, the dew point of the annealing atmosphere in the annealing process is preferably 30 °C or less from the viewpoint of effectively preventing oxidation of the surface of the Fe-based electroplated layer and improving coating adhesion when providing a galvanized layer.

(First cooling process)

The heated cold-rolled sheet is then cooled. The cooling involves a temperature range T₁ of T₂ or more and 750 °C or less. As a result of the cold-rolled sheet being cooled at the below-described first average cooling rate v₁ in the temperature range T₁, pearlite transformation is suppressed and bainite transformation can be utilized in the subsequent retention process.
If the first average cooling rate v₁ is excessively low, pearlite transformation occurs in the austenite generated in the first heating, untransformed austenite decreases, and bainite transformation cannot be utilized in the subsequent retention process. Hence, the area ratio of bainite decreases, and low-temperature toughness and stretch flangeability degrade. The first average cooling rate v₁ is therefore 2.0 °C/s or more, preferably 3.0 °C/s or more, and more preferably 5.0 °C/s or more.
Although no upper limit is placed on the first average cooling rate v₁, the first average cooling rate v₁ is preferably 60.0 °C/s or less from the viewpoint of reducing the burden of equipment investment.
Cooling in the temperature range T₁ is preferably continuous cooling.
The cooling rate from the first heating temperature to 750 °C is not limited.

(In-furnace retention process)

The cold-rolled sheet that has passed through the temperature range T₁ is then subjected to an in-furnace retention process in which it is held at a retention temperature T₂ of 350 °C or more and 550 °C or less. As a result of the cold-rolled sheet after the first cooling process being retained (held) at the retention temperature T₂ of 350 °C or more and 550 °C or less for a retention time t₂ (s) satisfying the condition that F defined in Formula 1 is 0.20 or more and 0.90 or less, bainite transformation occurs. By determining the retention time t₂ (s) from an expansion curve obtained from the Formaster test and applying it, the bainite transformation and the associated precipitation of iron carbides and distribution of C into untransformed austenite are optimized. Since the expansion curve depends on the steel composition and the thermal hysteresis up to the first cooling, it is necessary to determine the expansion curve for each steel composition and thermal hysteresis from the first heating temperature to T₂ °C and select an appropriate retention time t₂.
Formula 1 is as follows: $F = 1 - \exp (- {kt}^{n})$ where t is the retention time (s), and k and n are constants determined from the expansion curve of the Formaster test.
If the time (retention time t₂) during which the cold-rolled sheet is retained at the retention temperature T₂ is excessively short, bainite transformation is insufficient and the area ratio of bainite is low, and also the standard deviation of nanohardness is large, so that low-temperature toughness, stretch flangeability, and bendability of sheared end surface portions degrade. The retention time t₂ is therefore not less than time t with which F is 0.20, and preferably not less than time t with which F is 0.30.
If the time (retention time t₂) during which the cold-rolled sheet is retained at the retention temperature T₂ is excessively long, bainite transformation proceeds excessively, the amount of martensite decreases, and TS decreases. Moreover, cementite precipitates in untransformed austenite, hindering precipitation of metastable carbides in the subsequent second heating process. As a result, the proportion of the number of martensite blocks in which metastable carbides are present to the number of martensite blocks decreases, and low-temperature toughness degrades. The retention time t₂ is therefore not more than time t with which F is 0.90, and preferably not more than time t with which F is 0.80.
The relationship between F and t in Formula 1 is determined as follows.
The foregoing steel slab is subjected to the processes up to the first cooling, and then is retained at the retention temperature T₂ of 350 °C or more and 550 °C or less. The processes up to the first cooling are the hot rolling process, the pickling and cold rolling process, the first heating process, and the first cooling.
The expansion curve during retention at the retention temperature T₂ is obtained using a Formaster tester. Retention at T₂ is carried out until expansion stops. Setting the expansion amount at the start of retention at the retention temperature T₂ to 0 and the expansion amount at the time of stop to 1, the expansion curve is fitted with Formula 1 to calculate the constants k and n. This determines the relationship between F and t at the retention temperature T₂. $F = 1 - \exp (- {kt}^{n})$ where t is the retention time (s), and k and n are constants determined from the expansion curve of the Formaster test.

(Second cooling process)

Next, the cold-rolled sheet that has undergone the retention process is cooled. The cooling stop temperature is Ms - 20 °C or less. This allows martensitic transformation to proceed sufficiently. If the cooling stop temperature is more than Ms - 20 °C, untransformed austenite does not transform into martensite and the amount of retained austenite becomes excessive, making it impossible to achieve good component strength and stretch flangeability. The cooling stop temperature may be room temperature. Here, Ms is the temperature at which martensitic transformation begins to occur (Ms point), and the value measured by the below-described test is used.
During cooling, the cold-rolled sheet that has undergone the retention process passes through a temperature range T₃ of Ms - 20 °C or more and Ms °C or less. Here, by cooling the cold-rolled sheet in the temperature range T₃ at the below-described second average cooling rate v₂, the precipitation of cementite in martensite can be suppressed.
If the second average cooling rate v₂ is excessively low, cementite precipitates in martensite and the precipitation of metastable carbides in the subsequent second heating is suppressed, so that the proportion of the number of martensite blocks in which metastable carbides are present to the number of martensite blocks decreases and low-temperature toughness degrades.
The second average cooling rate v₂ is therefore 5 °C/s or more, and preferably 8 °C/s or more. Although no upper limit is placed on the second average cooling rate v₂, the second average cooling rate v₂ is preferably 60.0 °C/s or less from the viewpoint of reducing the burden of equipment investment.
The cooling rate outside the temperature range T₃ is not limited.
The Ms point is a value measured by the Formaster test as follows.
Using a Formaster tester, the foregoing steel slab is subjected to the processes up to the end of the retention process, and then cooled to room temperature at the second average cooling rate of 5 °C/s or more. The temperature at which martensitic transformation occurs and expansion begins during the second cooling is taken as the Ms point. No upper limit is placed on the second average cooling rate, and the second average cooling rate may be, for example, 100 °C/s or less.

(Second heating process)

The cold-rolled sheet cooled to the cooling stop temperature is then subjected to second heating.
As a result of the second heating, metastable carbides that improve low-temperature toughness are precipitated in the martensite blocks formed during the above-described cooling. Thus, a high-strength steel sheet having a TS of 780 MPa or more and excellent low-temperature toughness while ensuring good component strength, stretch flangeability, and bendability of sheared end surface portions is obtained.
The cooled cold-rolled sheet is heated to a second heating temperature X (unit: °C) and held for a holding time Y (unit: s) in the second heating. The temperature during the holding time Y is within the range of X ± 20 °C.
X and Y satisfy the following Formula 2. $7000 \leq (273 + X) \times (20 + \log (Y / 3600)) \leq 13000$
"(273 + X) × (20 + log(Y/3600))" in Formula 2 is hereafter referred to as "variable part Z" for convenience.
By performing the second heating that satisfies the foregoing conditions, the proportion of the number of martensite blocks in which metastable carbides are present to the number of martensite blocks (proportion p) can be increased to improve low-temperature toughness.
Here, the temperature X (°C) is higher than room temperature.
If the temperature X (°C) of the second heating is higher than the cooling stop temperature of the second cooling process, heating is performed from the cooling stop temperature to the temperature X (°C) of the second heating.
The cooling stop temperature of the second cooling process and the temperature X (°C) of the second heating may be the same. In this case, the cold-rolled sheet is held at the cooling stop temperature, which is the temperature X (°C) of the second heating.
If the value of the variable part Z in Formula 2 is excessively small, that is, if the temperature X is excessively low and/or the holding time Y is excessively short, metastable carbides do not precipitate sufficiently, so that the proportion p is low. The value of the variable part Z is therefore 7000 or more and preferably 8000 or more, from the viewpoint of increasing the proportion p.
If the value of the variable part Z is excessively high, that is, if the temperature X is excessively high and/or the holding time Y is excessively long, metastable carbides transition to cementite, so that the proportion p is low. The value of the variable part Z is therefore 13000 or less and preferably 12000 or less, from the viewpoint of increasing the proportion p.
In the second heating, the temperature X (unit: °C) preferably satisfies the following Formula 3. This increases the number density of metastable carbides in the martensite blocks in which metastable carbides are present (number density n). $100 \leq X \leq 400$
The temperature X is preferably 100 °C or more, more preferably 120 °C or more, and further preferably 150 °C or more, from the viewpoint of increasing the number density n.
From the same viewpoint, the temperature X is preferably 400 °C or less, more preferably 380 °C or less, and further preferably 350 °C or less.
The cold-rolled sheet that has undergone the second heating is then cooled to room temperature, for example. The cooling rate is not limited.
In this way, the high-strength steel sheet according to the present disclosure is obtained by the production method according to the present disclosure.
In the case where the production method according to the present disclosure involves the below-described coating or plating treatment, the obtained high-strength steel sheet according to the present disclosure is a coated or plated steel sheet including a coated or plated layer.
Other conditions of the series of heat treatments in the production method according to the present disclosure are not limited as long as the above-described thermal hysteresis is satisfied, and heat treatment lines, etc. are not limited.

(Second coating or plating process (optional))

In one embodiment of the present disclosure, the production method may include a second coating or plating process in which the steel sheet during the process from the first heating to the second heating is subjected to coating or plating treatment to form a second coated or plated layer, which is a metal coated or plated layer.
The second coating or plating process is preferably a galvanizing process or hot-dip aluminum coating treatment. Examples of galvanizing treatment in the galvanizing process include hot-dip galvanizing treatment, alloyed galvanizing treatment, electrogalvanizing treatment, and aluminum coating or plating treatment.
In the case of hot-dip galvanizing treatment, it is preferable to immerse the steel sheet in a galvanizing bath at 440 °C or more and 500 °C or less and then adjust the coating weight by gas wiping or the like. The hot-dip galvanizing bath is not limited as long as it has the composition of the galvanized layer described above, but it is preferable to use, for example, a molten bath having an Al content of 0.10 mass% or more and 0.23 mass% or less with the balance consisting of Zn and inevitable impurities.
In the case of alloyed galvanizing treatment, it is preferable to perform hot-dip galvanizing treatment in the above-described manner and then perform alloying treatment by heating the steel sheet including a hot-dip galvanized layer (hot-dip galvanized steel sheet) to an alloying temperature of 450 °C or more and 600 °C or less. If the alloying temperature is less than 450 °C, the Zn-Fe alloying rate is slow and alloying may be hindered. If the alloying temperature is more than 600 °C, untransformed austenite transforms into pearlite, making it difficult to achieve a TS of 590 MPa or more. The alloying temperature is more preferably 510 °C or more. The alloying temperature is more preferably 570 °C or less.
The coating weight of each of the steel sheet including a hot-dip galvanized layer (hot-dip galvanized steel sheet) (GI) and the steel sheet including a galvannealed layer (galvannealed steel sheet) (GA) is preferably 20 g/m² to 80 g/m² per side. The coating weight can be adjusted by gas wiping or the like.
In the case of performing hot-dip aluminum coating treatment which is another specific example of metal coating or plating treatment, the cold-rolled sheet obtained by the foregoing cold-rolled sheet annealing is immersed in an aluminum molten gas at 660 °C to 730 °C to perform hot-dip aluminum coating treatment, and then the coating weight is adjusted by gas wiping or the like.
In the case of performing electrogalvanizing treatment, the coating thickness is preferably in the range of 2 µm to 15 µm, without being limited thereto.

(Skin pass rolling (optional))

The obtained high-strength steel sheet may be subjected to skin pass rolling. Skin pass rolling may be performed after coating or plating treatment.
The rolling reduction ratio in skin pass rolling is preferably 0.05 % or more from the viewpoint of increasing yield stress. Although no upper limit is placed on the rolling reduction ratio, the upper limit is preferably 1.50 % from the viewpoint of productivity.
Skin pass rolling may be performed online or offline.
Skin pass with the desired rolling reduction ratio may be performed at once, or may be performed in several steps.
The foregoing series of treatments such as annealing and coating or plating is preferably performed in a continuous galvanizing line (CGL) from the viewpoint of productivity.

[Member and production method therefor]

A member according to the present disclosure and a production method therefor will be described.
The member according to the present disclosure is a member produced using the above-described high-strength steel sheet according to the present disclosure. For example, the member can be produced by forming the high-strength steel sheet according to the present disclosure into the desired shape by press working or the like.
The high-strength steel sheet according to the present disclosure is a high-strength steel sheet excellent in component strength, ductility, stretch flangeability, and bendability of sheared end surface portions. Hence, the use of the high-strength steel sheet according to the present disclosure or the member produced using the high-strength steel sheet for, for example, automotive framework structural components or automotive reinforcing components can make the automotive body more lightweight to improve fuel efficiency. The presently disclosed technology thus has very high industrial value.

EXAMPLES

The presently disclosed technology will be described in more detail below by way of examples, although the present disclosure is not limited to these examples.

[Test Nos. 1 to 51]

Molten steels having the chemical compositions shown in Table 1 below (with the balance consisting of Fe and inevitable impurities) were obtained by steelmaking in a converter, and continuously cast to obtain steel slabs.
The obtained steel slabs were each subjected to hot rolling to obtain a hot-rolled sheet. Specifically, the steel slab was heated to 1250 °C and rough-rolled under the conditions shown in Table 2 below, then finish-rolled at a finish rolling temperature of 900 °C, then coiled at 500 °C, and then cooled to room temperature to obtain a hot-rolled sheet.
The obtained hot-rolled sheet was pickled, then subjected to softening heat treatment at 500 °C, and then subjected to cold rolling with a rolling ratio of 50 %. A cold-rolled sheet with a sheet thickness of 1.6 mm was thus obtained.
The obtained cold-rolled sheet was heated to the first heating temperature shown in Table 2 and held for 200 s.
Following this, the cold-rolled sheet was cooled so that the average cooling rate in the temperature range T₁ (750 °C or less and T₂ or more) would be the first average cooling rate v₁ shown in Table 2, and subsequently cooled at the cooling rate to the retention temperature T₂ and retained at T₂ for the retention time t₂ shown in Table 2.
The parameter F and the retention time t₂ satisfy the following Formula 1. $F = 1 - \exp (- {kt}^{n})$ where t is the holding (retention) time (s), and k and n are constants determined from the expansion curve of the Formaster test.
k and n in Formula 1 were determined as follows.
The obtained cold-rolled sheet was cut into a width of 10 mm and a length of 3 mm, and subjected to the Formaster test. FTM-100 produced by Fuji Electronic Industrial Co., Ltd. was used for the Formaster test. In each example, the sample was heated to the first heating temperature shown in Table 2 and held for 200 s. Following this, the sample was cooled so that the average cooling rate in the temperature range T₁ (750 °C or less and T₂ or more) would be the first average cooling rate v₁ shown in Table 2, and subsequently cooled at the cooling rate to the retention temperature T₂ and retained at T₂ for 1000 s to obtain the expansion curve. For example, in Example 1, the sample was held at the first heating temperature of 800 °C for 200 s, then cooled to 480 °C at an average cooling rate of 18 °C/s, and then retained at 480 °C for 1000 s. The resultant expansion curve was fitted with Formula 1 to determine k and n.
Following the retention in a furnace at the retention temperature T₂ (350 °C or more and 550 °C or less), the cold-rolled sheet was cooled to the cooling stop temperature shown in Table 2 so that the average cooling rate in the temperature range T₃ (Ms - 20 °C or more and Ms °C or less) would be the second average cooling rate v₂ shown in Table 2.
The Ms point was determined as follows.
The obtained cold-rolled sheet was cut into a width of 10 mm and a length of 3 mm, and subjected to the Formaster test. FTM-100 produced by Fuji Electronic Industrial Co., Ltd. was used for the Formaster test. In each example, the sample was heated to the first heating temperature shown in Table 2 and held for 200 s. Following this, the sample was cooled so that the average cooling rate in the temperature range T₁ (750 °C or less and T₂ or more) would be the first average cooling rate v₁ shown in Table 2, subsequently cooled at the cooling rate to the retention temperature T₂ and retained at T₂ for t seconds, and then final-cooled to room temperature at 30 °C/s. The expansion curve during the final cooling was created, and the temperature at which expansion was observed was taken as the Ms point. For example, in Example 1, the sample was held at the first heating temperature of 800 °C for 200 s, then cooled to 480 °C at the first average cooling rate of 18 °C/s, then retained at 480 °C for 30 s, and then final-cooled to room temperature at 30 °C/s. The Ms point was determined from the resultant expansion curve.

(Coating or plating treatment)

Some cold-rolled sheets were, after being retained at the retention temperature T₂, subjected to hot-dip galvanizing treatment to form a coated layer (hot-dip galvanized layer) on both sides. Thus, a hot-dip galvanized steel sheet (GI) was obtained. The steel sheet was then cooled to the cooling stop temperature at the second average cooling rate v₂.
For the hot-dip galvanizing treatment, a hot-dip galvanizing bath (bath temperature: 470 °C) containing Al: 0.20 mass% with the balance consisting of Zn and inevitable impurities was used.
The coating weight of the hot-dip galvanized layer per side was about 45 g/m² to 72 g/m².
The composition of the formed hot-dip galvanized layer contained Fe: 0.1 mass% to 1.0 mass% and Al: 0.2 mass% to 1.0 mass% with the balance consisting of Fe and inevitable impurities.
Some other cold-rolled sheets were, after being retained at the retention temperature T₂, subjected to galvannealing treatment to form a coated layer (galvannealed layer) on both sides. Thus, a galvannealed steel sheet (GA) was obtained. The steel sheet was then cooled to the cooling stop temperature at the second average cooling rate v₂.
For the hot-dip galvanizing treatment, a hot-dip galvanizing bath (bath temperature: 470 °C) containing Al: 0.14 mass% with the balance consisting of Zn and inevitable impurities was used.
The alloying treatment temperature was 550 °C.
The coating weight of the galvannealed layer per side was about 45 g/m².
The composition of the formed galvannealed layer contained Fe: 7 mass% to 15 mass% and Al: 0.1 mass% to 1.0 mass% with the balance consisting of Fe and inevitable impurities.
In the "Coating type" column in Table 3, "GI" indicates that a hot-dip galvanized layer was formed, "GA" indicates that a galvannealed layer was formed, and "CR" indicates that no coated layer was formed.
The steel sheet cooled to the cooling stop point was reheated to the temperature X [°C] shown in Table 2 and held for the holding time Y [s]. The variable part Z is (273 + X) × (20 + log(Y/3600)).

< Observation of steel microstructure>

For each of the obtained steel sheets, the area ratios of martensite, bainite, ferrite, and retained austenite, the proportion of the number of martensite blocks in which metastable carbides were present to the number of martensite blocks (proportion p), the average number density of metastable carbides in the martensite blocks in which metastable carbides were present (number density n), and the standard deviation of nanohardness (225 measurement points) were measured. The results are shown in Table 3 below.
The area ratio of residual microstructure was also measured by a commonly known method. Regarding the residual microstructure, "θ" in Table 3 means cementite precipitated in ferrite.

Each of the obtained steel sheets was subjected to the below-described tests to evaluate various properties. The results are shown in Table 3.

(Tensile test)

A tensile test was conducted in accordance with JIS Z 2241.
In detail, a JIS No. 5 test piece was collected from the obtained steel sheet so that the direction orthogonal to the rolling direction of the steel sheet would be the longitudinal direction. A tensile test was conducted using the collected test piece at a crosshead speed of 1.67 × 10^-1 mm/s to measure yield stress (YS) [MPa], tensile strength (TS) [MPa], and total elongation (El) [%]. Further, the yield ratio (YR) (= 100 × YS/TS) [%] was calculated.
If the tensile strength (TS) was 780 MPa or more, the steel sheet was determined to have high strength.
If the yield ratio (YR) was 55 % or more, the steel sheet was determined to have excellent component strength.
If the total elongation (El) was 10 % or more, the steel sheet was determined to have excellent ductility.

(Hole expanding test)

A hole expanding test was conducted in accordance with JIS Z 2256.
In detail, the obtained steel sheet was sheared to collect a test piece of 100 mm × 100 mm in size, and a hole with a diameter of 10 mm was drilled through the collected test piece with clearance 12.5 %. Next, in a state in which the test piece was clamped using a die having an inner diameter of 75 mm with a blank holding force of 9 tons (88.26 kN), a conical punch with an apical angle of 60° was pushed into the hole and the hole diameter Df [mm] at crack initiation limit was measured. The hole expansion ratio λ [%] was calculated using the following formula. $λ = \{(Df - D 0) / D 0\} \times 100$ where D0 [mm] is the initial hole diameter.
If the hole expansion ratio (λ) was 20 % or more, the steel sheet was determined to have excellent stretch flangeability.

(Sheared end surface portion bend test)

A bend test was conducted in accordance with JIS Z 2248.
In detail, a strip test piece of 30 mm in width and 100 mm in length was collected from the obtained steel sheet so that the axial direction of the bend test would be parallel to the rolling direction of the steel sheet. In the bend test of the sheared end surface sample, the end surface in the longitudinal direction was the sheared end surface. In the bend test of the ground end surface sample, the end surface in the longitudinal direction was the ground end surface.
Using the collected test piece, a 90° V bend test was conducted with an indentation load of 100 kN and a press holding time of 5 seconds.
The bend test was conducted on five test pieces with an appropriate bending radius R. Next, whether cracks occurred in the ridge line part of the bending apex was determined.
Whether cracks occurred was determined by observing the ridge line part of the bending apex with a digital microscope (RH-2000 produced by HIROX Co., Ltd.) at 40x magnification.
The minimum bending radius R at which no cracks occurred in any of the five test pieces was determined, and the value (R/t) obtained by dividing the minimum bending radius R by the sheet thickness t was taken as the limit bending radius. The limit bending radius (Rs/t) of the sheared end surface sample and the limit bending radius (Rg/t) of the ground end surface sample were calculated. If the ratio (Rs/Rg) of the limit bending radius (Rs/t) of the sheared end surface sample to the limit bending radius (Rg/t) of the ground end surface sample was 1.50 or less, the steel sheet was determined to have excellent bendability of sheared end surface portions.

(Low-temperature toughness test)

Six steel sheets (steel sheets with a thickness t of 1.6 mm) obtained were stacked and bonded together to prepare a Charpy impact test piece with a thickness t of 9.6 mm. The notch was a 2 mm U-notch. Using the test piece, a Charpy impact test was conducted at -40 °C, and the resultant Charpy absorbed energy was measured. In addition, the fracture surface after the test was observed to measure the percent ductile fracture.
The product of the Charpy absorbed energy (unit: J) and the percent ductile fracture (%) (Charpy absorbed energy (J) × percent ductile fracture (%)) was calculated and defined as the low-temperature toughness parameter P.
If the low-temperature toughness parameter P was 3000 or more, the steel sheet was determined to have excellent low-temperature toughness.

[Table 2]

Table 2

No.	Steel sample ID	Average strain rate of rough rolling [s^-1]	Total rolling reduction ratio of rough rolling (%)	First heating temperature [°C]	First average cooling rate v1 (temperature range T1) [°C/s]	Retention temperature T2 [°C]	Parameter F	Retention time t2 (temperature T2) [s]	Second average cooling rate v2 (temperature range T3) [°C/s]	Cooling stop temperature [°C]	Temperature X [°C]	Holding time Y [s]	Variable part Z	Ms point [°C]	Coating type	Remarks
1	A	0.0060	72	800	18	480	0.62	30	30	50	200	65	8635	358	GA	Example
2	A	0.0600	72	810	15	470	0.63	34	30	50	220	54	8961	341	GA	Example
3	A	0.6000	70	810	15	475	0.65	40	30	50	200	65	8635	338	GA	Comparative Example
4	A	0.0006	70	800	20	480	0.47	25	30	50	190	80	8495	355	GA	Example
5	A	0.0001	70	810	20	485	0.46	25	30	50	200	65	8635	352	GA	Comparative Example
6	A	0.0060	55	800	20	480	0.66	42	30	50	200	60	8619	359	GI	Example
7	A	0.0060	40	800	18	480	0.67	45	30	50	200	60	8619	361	GA	Comparative Example
8	A	0.0060	72	740	20	460	0.69	72	30	50	210	60	8801	73	GA	Comparative Example
9	A	0.0080	72	810	2.7	490	0.72	87	30	50	210	65	8818	360	CR	Example
10	A	0.0060	72	810	1.6	380	0.73	89	30	50	210	65	8818	354	GA	Comparative Example
11	A	0.0080	72	800	20	480	0.27	11	30	50	200	65	8635	352	CR	Example
12	A	0.0060	76	800	20	480	0.14	7	30	50	200	65	8635	357	CR	Comparative Example
13	A	0.0080	76	800	20	480	0.82	97	30	50	220	65	9001	200	GI	Example
14	A	0.0060	80	810	18	480	0.96	127	30	50	200	65	8635	26	GA	Comparative Example
15	A	0.0060	80	790	18	495	0.62	30	3	50	200	65	8635	358	GA	Comparative Example
16	A	0.0080	72	790	20	490	0.62	30	8	400	200	65	8635	367	GA	Comparative Example
17	A	0.0080	75	800	18	485	0.62	30	8	25	95	20000	7634	355	GA	Example
18	A	0.0080	72	800	15	480	0.62	30	5	25	420	30	12419	355	GA	Example
19	A	0.0080	78	800	15	475	0.62	30	35	120	110	5	6566	354	GA	Comparative Example
20	A	0.0060	72	800	18	480	0.62	30	4	20	380	8000	13286	344	GA	Comparative Example
21	B	0.0080	69	810	18	485	0.72	28	10	300	300	30	10269	351	GA	Example
22	C	0.0060	69	830	18	480	0.47	32	50	50	220	65	9001	346	GA	Example
23	D	0.0060	69	800	20	480	0.76	28	100	50	200	65	8635	294	GI	Example
24	E	0.0080	72	800	20	480	0.65	34	100	50	200	65	8635	315	CR	Example
25	F	0.0060	72	770	20	480	0.60	30	100	50	250	30	9373	332	GA	Example
26	G	0.0060	72	800	20	480	0.72	15	150	200	300	25	10223	156	GA	Example
27	H	0.0060	70	800	20	480	0.60	100	200	50	250	65	9548	303	GA	Example
28	I	0.0060	70	800	18	485	0.60	15	300	50	205	65	8727	326	GA	Comparative Example
29	J	0.0080	70	850	18	480	0.60	35	30	50	220	65	9001	340	GA	Comparative Example
30	K	0.0060	70	800	18	480	0.60	28	30	50	180	65	8270	210	GA	Comparative Example
31	L	0.0060	70	820	18	520	0.60	125	30	50	200	65	8635	272	GA	Comparative Example
32	M	0.0080	70	800	18	430	0.60	50	30	50	250	65	9548	355	GA	Example
33	N	0.0060	72	850	14	380	0.60	100	30	200	300	20	10168	332	GA	Example
34	O	0.0060	72	800	18	480	0.60	24	30	50	180	65	8270	339	GA	Example
35	P	0.0060	72	795	18	480	0.60	55	30	50	200	65	8635	351	GA	Example
36	Q	0.0060	72	800	18	480	0.60	35	30	300	300	60	10441	359	GI	Example
37	R	0.0080	72	830	18	480	0.55	45	10	50	300	65	10461	356	GA	Example
38	S	0.0060	72	800	18	480	0.60	35	20	50	200	65	8635	344	GA	Example
39	T	0.0060	72	800	20	450	0.65	40	30	40	200	65	8635	364	CR	Example
40	U	0.0080	68	820	20	480	0.60	30	30	30	200	65	8635	300	GA	Example
41	V	0.0060	68	770	25	450	0.67	35	10	50	180	200	8491	357	GA	Example
42	W	0.0080	65	800	30	480	0.60	35	30	60	200	65	8635	337	CR	Example
43	X	0.0060	72	800	30	480	0.45	28	30	50	200	65	8635	339	GA	Example
44	Y	0.0060	69	830	30	420	0.56	40	9	40	200	65	8635	350	GA	Example
45	Z	0.0060	69	800	18	480	0.45	20	30	50	140	1000	8030	347	GI	Example
46	AA	0.0060	72	810	40	480	0.35	19	30	200	200	65	8635	357	GA	Example
47	AB	0.0060	72	810	18	450	0.55	32	30	25	180	300	8571	346	GA	Example
48	AC	0.0080	72	810	18	470	0.75	100	30	50	200	65	8635	351	GA	Example
49	AD	0.0060	70	800	18	470	0.60	30	8	200	200	30	8477	362	GA	Example
50	AE	0.0080	70	800	18	485	0.60	30	30	25	150	12000	8681	361	GA	Example
51	AF	0.0060	72	800	21	490	0.60	30	30	50	220	65	9001	350	CR	Example

Underlines indicate outside the range according to the present disclosure.

[Table 3]

Table 3

No.	Steel sample ID	Area ratio [%]					Proportion p [%]	Number density n [10⁶/mm²]	Average nanohardness [Hn]^ave	σ_n/[Hn]^ave*	YS [MPa]	TS [MPa]	YR [%]	El [%]	[%]	Sheared end surface portion bendability [-]	Low-temperature toughness	Remarks
No.	Steel sample ID	Martensite	Bainite	Ferrite	Retailed austenite	Residual microstructure	Proportion p [%]	Number density n [10⁶/mm²]	Average nanohardness [Hn]^ave	σ_n/[Hn]^ave*	YS [MPa]	TS [MPa]	YR [%]	El [%]	[%]	Sheared end surface portion bendability [-]	Low-temperature toughness	Remarks
1	A	55	5	37	2	1	64	600	5.2	0.31	702	990	71	17	31	1.21	5984	Example
2	A	42	5	51	0	2	56	300	4.8	0.52	732	996	73	15	27	1.40	5258	Example
3	A	40	5	54	0	1	50	400	4.8	0.64	685	1018	67	16	32	1.60	4898	Comparative Example
4	A	52	7	37	2	2	56	400	5.0	0.51	651	992	66	17	25	1.40	5833	Example
5	A	49	8	39	2	2	48	300	5.2	0.62	667	993	67	16	29	1.60	5152	Comparative Example
6	A	59	7	33	0	1	48	400	5.3	0.51	731	991	74	17	25	1.40	5190	Example
7	A	60	8	28	2	2	48	500	5.4	0.63	735	996	74	17	25	1.60	5451	Comparative Example
8	A	8	6	85	0	1	48	200	4.3	0.37	502	768	65	22	34	1.21	5101	Comparative Example
9	A	60	3	34	0	3	48	400	4.7	0.37	657	1001	66	15	22	1.21	3069	Example
10	A	52	1	45	1	1	30	200	5.3	0.40	736	999	74	15	31	1.21	2789	Comparative Example
11	A	49	3	44	2	2	30	200	4.9	0.51	702	1019	69	16	23	1.40	3135	Example
12	A	54	1	41	2	2	30	200	4.8	0.68	681	1005	68	15	15	1.60	2758	Comparative Example
13	A	12	25	60	1	2	4	5	5.0	0.32	630	981	64	15	28	1.21	3321	Example
14	A	7	29	62	0	2	0	0	4.7	0.37	620	962	64	17	25	1.21	2684	Comparative Example
15	A	57	4	38	0	1	0	0	5.0	0.39	720	1026	70	16	31	1.21	2899	Comparative Example
16	A	55	6	21	16	2	36	25	5.1	0.32	500	1003	50	16	12	1.21	4712	Comparative Example
17	A	53	4	41	1	1	4	0.8	5.3	0.32	701	1006	70	16	33	1.21	3247	Example
18	A	52	6	39	2	1	20	0.6	5.5	0.30	632	1022	62	15	25	1.21	3019	Example
19	A	51	7	39	2	1	0	0	4.8	0.30	643	1026	63	17	29	1.21	2952	Comparative Example
20	A	44	6	48	0	2	0	0	5.1	0.39	709	994	71	15	34	1.21	2475	Comparative Example
21	B	29	5	62	2	2	16	8	4.1	0.47	532	825	64	19	35	1.25	5942	Example
22	C	60	32	5	1	2	60	100	6.6	0.36	835	1214	69	15	34	1.21	5442	Example
23	D	14	8	75	1	2	50	100	4.2	0.33	542	802	68	20	35	1.25	5282	Example
24	E	78	8	0	12	2	60	200	8.3	0.35	895	1538	58	11	23	1.21	5793	Example
25	F	28	8	52	10	2	60	58	5.4	0.32	569	1011	56	17	22	1.21	5697	Example
26	G	13	7	77	1	2	24	12	3.8	0.35	501	789	63	22	26	1.25	5959	Example
27	H	25	6	53	14	2	24	400	5.2	0.37	632	1121	56	17	21	1.21	5568	Example
28	I	8	4	84	2	2	20	50	2.3	0.38	298	437	68	29	34	1.21	5570	Comparative Example
29	J	12	6	64	17	1	20	25	5.3	0.31	503	992	51	19	17	1.21	5009	Comparative Example
30	K	7	8	82	2	1	20	250	3.9	0.37	402	697	58	28	28	1.21	5098	Comparative Example
31	L	28	4	46	20	2	50	90	5.4	0.34	498	1020	49	15	12	1.21	5695	Comparative Example
32	M	51	6	40	2	1	32	500	5.6	0.40	717	1024	70	16	30	1.21	5834	Example
33	N	59	30	8	1	2	24	36	6.5	0.36	845	1214	70	14	32	1.21	5442	Example
34	O	25	4	69	1	1	50	400	4.3	0.40	562	832	68	18	32	1.25	5455	Example
35	P	43	5	51	0	1	60	400	4.9	0.31	637	993	64	17	27	1.21	5376	Example
36	Q	42	5	49	2	2	30	10	5.4	0.40	711	1017	70	16	34	1.21	5745	Example
37	R	62	28	7	2	1	26	12	5.8	0.33	850	1198	71	13	28	1.21	5295	Example
38	S	40	4	52	2	2	70	800	5.4	0.32	641	1023	63	15	33	1.21	5736	Example
39	T	58	4	34	2	2	60	600	5.0	0.37	642	999	64	15	33	1.21	5999	Example
40	U	36	62	0	1	1	60	620	6.3	0.40	854	1234	69	12	36	1.21	5031	Example
41	V	31	7	59	2	1	40	300	4.0	0.34	520	824	63	19	35	1.25	5661	Example
42	W	43	5	50	1	1	60	300	5.4	0.30	639	1011	63	16	30	1.21	5462	Example
43	X	40	4	55	0	1	60	300	5.5	0.39	666	996	67	17	31	1.21	5099	Example
44	Y	62	32	4	1	1	50	470	5.9	0.35	902	1214	62	12	29	1.21	5815	Example
45	Z	44	6	47	2	1	50	500	5.4	0.33	622	1016	61	15	28	1.21	5194	Example
46	AA	54	7	37	1	1	50	300	4.8	0.36	652	997	65	17	32	1.21	5828	Example
47	AB	44	8	45	2	1	40	400	5.7	0.30	716	1017	70	16	33	1.21	5949	Example
48	AC	46	5	46	2	1	40	470	4.8	0.31	699	1019	69	16	33	1.21	5363	Example
49	AD	31	8	58	2	1	40	400	4.0	0.36	447	834	67	21	30	1.50	5671	Example
50	AE	49	6	42	2	1	70	300	5.3	0.36	677	1000	68	15	33	1.21	5661	Example
51	AF	46	5	45	3	1	40	470	5.2	0.36	708	997	71	16	33	1.21	5088	Example

*σ_n is standard deviation of nanohardness.
Underliies ildicate outside the range according to the present disclosure.

As shown in Table 3, all Examples had a TS of 780 MPa or more and were excellent in component strength, ductility, stretch flangeability, bendability of sheared end surface portions, and low-temperature toughness. In Comparative Examples, on the other hand, at least one out of component strength, ductility, stretch flangeability, bendability of sheared end surface portions, and low-temperature toughness was poor.

[Test Nos. 52 to 75]

Steel slabs (steel materials) having the chemical compositions shown in Table 1 with the balance consisting of Fe and inevitable impurities were obtained by steelmaking using a converter and continuous casting. The obtained steel slabs were each heated to 1250 °C and subjected to rough rolling to obtain a sheet bar. The obtained sheet bar was then subjected to finish rolling at a finish rolling temperature of 900 °C and coiled under the conditions shown in Table 4 below to obtain a hot-rolled sheet. The obtained hot-rolled sheet was pickled and then subjected to cold rolling under the conditions shown in Table 4 to obtain a cold-rolled sheet with a sheet thickness of 1.2 mm.
The obtained cold-rolled steel was subjected to the first coating or plating process (metal electroplating process), the first heating process, the first cooling process, the retention in a furnace process, the second coating or plating process, the second cooling process, and the second heating process under the conditions shown in Table 4 to obtain a steel sheet.
In the "Metal electroplating treatment (plating type)" column in Table 4, "Applied (Fe)" is an example where Fe-based electroplating treatment was applied, and "Applied (Ni)" is an example where Ni-based electroplating treatment was applied. The composition of the metal electroplated layer by Fe-based electroplating contained Fe: 95 mass% to 100 mass% with the balance consisting of inevitable impurities. The composition of the metal electroplated layer by Ni-based electroplating contained Ni: 95 mass% to 100 mass% with the balance consisting of inevitable impurities
In the second coating or plating process, some hot-rolled sheets (pickled skin) or cold-rolled sheets were subjected to hot-dip galvanizing treatment or galvannealing treatment to obtain hot-dip galvanized steel sheets (hereafter also referred to as GI) or galvannealed steel sheets (hereafter also referred to as GA).
The galvanizing bath temperature was 470 °C for both GI and GA production.
The galvanizing coating weight was 45 g/m² to 75 g/m² per side (double-sided coating) for GI production, and 40 g/m² to 65 g/m² per side (double-sided coating) for GA production.
The composition of the galvanized layer of the finally obtained steel sheet contained Fe: 0.1 mass% to 1.0 mass% and Al: 0.20 mass% to 0.33 mass% with the balance consisting of Zn and inevitable impurities in the case of GI, and contained Fe: 7.0 mass% to 12.0 mass% and Al: 0.10 mass% to 0.23 mass% with the balance consisting of Zn and inevitable impurities in the case of GA.
Each of the hot-dip galvanized layer and the galvannealed layer was formed on both sides of the base steel sheet.

[Thickness of soft surface layer]

The soft surface layer was measured as follows. After peeling off the coated or plated layer (hot-dip galvanized layer or alloyed galvanized layer, or metal electroplated layer in some cases), a sheet thickness section (L-section) parallel to the rolling direction of the base steel sheet was smoothed by wet polishing, and then measurement was performed using a Vickers hardness meter with a load of 10 gf from a position of 1 µm in the sheet thickness direction from the surface of the steel sheet to a position of 100 µm in the sheet thickness direction at 1 µm intervals. After this, measurement was performed at 20 µm intervals up to the sheet thickness center. A region where the hardness was 85 % or less of the hardness at the position of 1/4 of the sheet thickness was defined as the soft layer (soft surface layer), and the thickness of the region in the sheet thickness direction was taken as the thickness of the soft layer.

[Nanohardness of soft surface layer]

After peeling off the coated or plated layer (hot-dip galvanized layer or alloyed galvanized layer, or metal electroplated layer in some cases), mechanical polishing, buffing with diamond and alumina, and colloidal silica polishing were performed from the surface of the base steel sheet to the position of 1/4 of the depth of the soft surface layer from the surface of the base steel sheet in the sheet thickness direction. A nanoindentation device (Tribo-950 produced by Hysitron, Inc.) was used to measure the nanohardness at a total of 512 points with a Berkovich diamond indenter under the following conditions:

Loading and unloading speed: 50 µN/s
Maximum load: 500 µN
Measurement region: 50 µm × 50 µm
Data collection pitch: 5 msec
Indentation spacing: 2 µm.

Next, mechanical polishing, buffing with diamond and alumina, and colloidal silica polishing were performed to the position of 1/2 of the depth of the soft surface layer in the sheet thickness direction. Tribo-950 produced by Hysitron, Inc. was used to measure the nanohardness at a total of 512 points with a Berkovich diamond indenter under the same conditions as above.
For each of the obtained steel sheets, component strength, ductility, stretch flangeability, bendability of sheared end surface portions, and low-temperature toughness were evaluated according to the foregoing test methods. The results are shown in Table 5 below.
Moreover, a V bend + orthogonal VDA bend test and an axial crush fracture test were conducted according to the following test methods. The results are shown in Table 5.
In the case of hot-dip galvanized steel sheets with a sheet thickness of more than 1.2 mm, the V bend + orthogonal VDA bend test and the axial crush test were conducted on steel sheets with a sheet thickness of 1.2 mm, taking into account the influence of the sheet thickness. Each steel sheet with a sheet thickness of more than 1.2 mm was ground on one side to make the sheet thickness 1.2 mm. In the case of hot-dip galvanized steel sheets with a sheet thickness of less than 1.2 mm, the V bend + orthogonal VDA bend test and the axial crush test were conducted without grinding because the influence of the sheet thickness was insignificant.

(V bend + orthogonal VDA bend test)

The V bend + orthogonal VDA bend test was conducted as follows.
A 60 mm × 65 mm test piece was collected from the obtained steel sheet by shearing and end grinding, with the 60 mm side parallel to the rolling (L) direction. A test piece was prepared by 90° bending (primary bending) in the rolling (L) direction with the width (C) direction as the axis at a curvature radius/sheet thickness of 4.2. In the 90° bending (primary bending), a punch B1 was pressed against the steel sheet placed on a die A1 having a V groove to obtain a test piece T1, as illustrated in FIG. 2A. Next, a punch B2 was pressed against the test piece T1 placed on support rolls A2 so that the bending direction would be orthogonal to the rolling direction to perform orthogonal bending (secondary bending), as illustrated in FIG. 2B. In FIGS. 2A and 2B, D1 indicates the width (C) direction and D2 indicates the rolling (L) direction.
The V bending conditions in the V bend + orthogonal VDA bend test are as follows:

Test method: die support, punch press
Forming load: 10 tons
Test speed: 30 mm/min
Holding time: 5 seconds
Bending direction: rolling (L) direction.

The VDA bending conditions in the V bend + orthogonal VDA bend test are as follows:

Test method: roll support, punch press
Roll diameter: ϕ30 mm
Punch tip R: 0.4 mm
Roll distance: (sheet thickness × 2) + 0.5 mm
Stroke speed: 20 mm/min
Test piece size: 60 mm × 60 mm
Bending direction: direction (C) orthogonal to rolling direction.

The stroke at maximum load was determined in the stroke-load curve obtained when the VDA bending was performed. The average value of the stroke at maximum load when the V bend + orthogonal VDA bend test was conducted three times was taken as SF_max (mm). Collision fracture resistance (resistance to bending fracture) was determined as excellent if SF_max was 26.0 mm or more.

(Axial crush test)

The axial crush test was conducted as follows.
A 150 mm × 100 mm test piece was collected from the obtained steel sheet by shearing, with the 150 mm side parallel to the rolling (L) direction. Using a die with a punch shoulder radius of 5.0 mm and a die shoulder radius of 5.0 mm, a hat-shaped member 10 illustrated in FIGS. 3A and 3B was prepared by forming (bending) to a depth of 40 mm.
Moreover, the steel sheet used as the material for the hat-shaped member was cut out to a size of 80 mm × 100 mm. A cut-out steel sheet 20 and the hat-shaped member 10 were spot welded to prepare a test member 30 illustrated in FIGS. 3A and 3B. FIG. 3A is a front view of the test member 30 prepared by spot welding the hat-shaped member 10 and the steel sheet 20. FIG. 3B is a perspective view of the test member 30. The positions of spot weld portions 40 were such that the distance between the end of the steel sheet and the weld was 10 mm and the distance between the welds was 20 mm, as illustrated in FIG. 3B. Next, as illustrated in FIG. 3C, the test member 30 was joined to a base plate 50 by TIG welding to prepare an axial crush test sample. An impactor 60 was then made to collide with the prepared axial crush test sample at a constant collision velocity of 10 mm/min to crush the axial crush test sample by 70 mm. The crush direction D3 was parallel to the longitudinal direction of the test member 30, as illustrated in FIG. 3C.
The appearance of the test model 30 after the test was observed to check whether there was axial crush fracture (appearance crack).
In Table 5, "A" indicates that no appearance crack was found, "B" indicates that an appearance crack was found in one location, and "C" indicates that an appearance crack was found in two or more locations. The collision fracture resistance (resistance to axial crush fracture) was determined as excellent in the case of "A" or "B".

[Table 4]

Table 4

No.	Steel sample ID	Average strain rate of rough rolling [s^-1]	Total rolling reduction ratio of rough rolling [%]	Metal electroplating treatment (plating type)	First heating temperature [°C]	Dew point of first heating temperature [°C]	First average cooling rate v1 (temperature range T1) [°C/s]	Parameter F	Retention time t2 (temperatureT2) [s]	Second average cooling rate v2 (temperature range T3) [°C/s]	Cooling stop temperature [°C]	Temperature X [°C]	Holding time Y [s]	Variable part Z	Ms point [°C]	Coating type	Remarks
52	A	0.003	70	Not applied	810	-20	14	0.62	35	24	30	200	50	8581	355	GA	Example
53	A	0.006	72	Not applied	810	5	15	0.62	35	24	30	200	50	8581	356	GA	Example
54	A	0.006	70	Applied (Fe)	810	-20	16	0.62	35	25	30	200	50	8581	355	GA	Example
55	A	0.004	72	Applied (Fe)	810	5	14	0.62	35	23	30	200	50	8581	357	GA	Example
56	A	0.003	69	Applied (Ni)	810	5	15	0.62	35	25	30	200	50	8581	355	CR	Example
57	A	0.008	72	Not applied	810	5	16	0.62	35	23	30	200	50	8581	361	GI	Example
58	A	0.002	70	Applied (Fe)	810	-20	14	0.62	35	25	30	200	50	8581	360	GI	Example
59	A	0.004	69	Applied (Fe)	810	5	17	0.62	35	23	30	200	50	8581	356	GI	Example
60	B	0.008	66	Not applied	790	-25	17	0.72	28	15	350	300	100	10568	354	GA	Example
61	B	0.006	66	Not applied	790	-5	16	0.72	28	14	350	300	100	10568	349	GA	Example
62	B	0.009	67	Applied (Fe)	790	-25	16	0.72	28	14	350	300	100	10568	349	GA	Example
63	B	0.007	66	Applied (Fe)	790	-5	18	0.72	28	17	350	300	100	10568	347	GA	Example
64	B	0.007	67	Applied (Ni)	790	-5	17	0.72	28	17	350	300	100	10568	349	GA	Example
65	B	0.003	66	Not applied	790	-5	16	0.72	28	16	350	300	100	10568	344	GI	Example
66	B	0.003	66	Applied (Fe)	790	-25	17	0.72	28	14	350	300	100	10568	356	GI	Example
67	B	0.008	65	Applied (Fe)	790	-5	18	0.72	28	16	350	300	100	10568	349	GI	Example
68	R	0.007	74	Not applied	830	-15	14	0.60	40	20	50	300	40	10340	356	GA	Example
69	R	0.002	74	Not applied	830	10	15	0.60	40	20	50	300	40	10340	357	GA	Example
70	R	0.009	72	Applied (Fe)	830	-15	14	0.60	40	19	50	300	40	10340	359	GA	Example
71	R	0.002	73	Applied (Fe)	830	10	13	0.60	40	19	50	300	40	10340	357	GA	Example
72	R	0.007	74	Applied (Ni)	830	10	14	0.60	40	22	50	300	40	10340	358	GA	Example
73	R	0.009	74	Not applied	830	-5	16	0.60	40	22	50	300	40	10340	357	GI	Example
74	R	0.005	74	Applied (Fe)	830	-25	16	0.60	40	19	50	300	40	10340	357	GI	Example
75	R	0.002	72	Applied (Fe)	830	-5	15	0.60	40	22	50	300	40	10340	357	GI	Example

Underlines indicate outside the range according to the present disclosure.

As shown in Table 5, Examples had a TS of 780 MPa or more and were excellent in component strength, ductility, stretch flangeability, bendability of sheared end surface portions, low-temperature toughness, and collision fracture resistance (bending fracture resistance and axial crush resistance).
It is thus possible to provide a high-strength steel sheet excellent in component strength, stretch flangeability, bendability of sheared end surface portions, low-temperature toughness, and collision fracture resistance (bending fracture resistance and axial crush resistance). The high-strength steel sheet has various excellent properties, and accordingly can be used for automotive framework structural components and reinforcing components of various sizes and shapes, etc. This can make the automotive body more lightweight to improve fuel efficiency. The presently disclosed technology thus has very high industrial value.

REFERENCE SIGNS LIST

A1: die
A2: support roll
B1: punch
B2: punch
T1: test piece
D1: width (C) direction
D2: rolling (L) direction
D3: crushing direction
10: hat-shaped member
20: steel sheet
30: test member
40: spot weld portion
50: base plate
60: impactor

Claims

A high-strength steel sheet comprising:
a chemical composition containing, in mass%,

C: 0.030 % or more and 0.500 % or less,

Si: 0.01 % or more and 2.50 % or less,

Mn: 0.10 % or more and 5.00 % or less,

P: 0.100 % or less,

S: 0.0200 % or less,

Al: 1.000 % or less,

N: 0.0100 % or less, and

O: 0.0100 % or less

with a balance consisting of Fe and inevitable impurities; and

a steel microstructure at a position of 1/4 of a sheet thickness of the high-strength steel sheet in which an area ratio of martensite is 10 % or more and 80 % or less, an area ratio of bainite is 2 % or more and 70 % or less, an area ratio of ferrite is 80 % or less, an area ratio of retained austenite is 15 % or less, and a proportion of a number of martensite blocks in which metastable carbides are present to a total number of martensite blocks is 2 % or more,

wherein when nanohardness is measured at 225 points or more at the position of 1/4 of the sheet thickness, a standard deviation σ_n of the nanohardness is 0.60 × [H_n]^ave or less, where [H_n]^ave is an average value of the nanohardness.
The high-strength steel sheet according to claim 1, wherein an average number density of the metastable carbides in the martensite blocks in which the metastable carbides are present is 1 × 10⁶/mm² or more.
The high-strength steel sheet according to claim 1 or 2, wherein the chemical composition further contains, in mass%, at least one element selected from the group consisting of
Ti: 0.200 % or less,

Nb: 0.200 % or less,

V: 0.200 % or less,

Ta: 0.10 % or less,

W: 0.10 % or less,

B: 0.0100 % or less,

Cr: 1.00 % or less,

Mo: 1.00 % or less,

Ni: 1.00 % or less,

Co: 0.010 % or less,

Cu: 1.00 % or less,

Sn: 0.200 % or less,

Sb: 0.200 % or less,

Ca: 0.0100 % or less,

Mg: 0.0100 % or less,

REM: 0.0100 % or less,

Zr: 0.100 % or less,

Te: 0.100 % or less,

Hf: 0.10 % or less, and

Bi: 0.200 % or less.
The high-strength steel sheet according to any one of claims 1 to 3, comprising a soft surface layer that is a region of 200 µm or less from a surface of the high-strength steel sheet in a sheet thickness direction and having a Vickers hardness of 85 % or less of a Vickers hardness at the position of 1/4 of the sheet thickness of the high-strength steel sheet,
wherein when nanohardness is measured at 300 points or more in a 50 µm × 50 µm region of a sheet plane at each of a position of 1/4 of a depth of the soft surface layer from the surface of the high-strength steel sheet in the sheet thickness direction and a position of 1/2 of the depth of the soft surface layer from the surface of the high-strength steel sheet in the sheet thickness direction,

a proportion of a number of measurements in which the nanohardness of the sheet plane at the position of 1/4 of the depth of the soft surface layer from the surface of the high-strength steel sheet in the sheet thickness direction is 7.0 GPa or more is 0.10 or less relative to a total number of measurements,

a standard deviation σ of the nanohardness of the sheet plane at the position of 1/4 of the depth of the soft surface layer from the surface of the high-strength steel sheet in the sheet thickness direction is 1.8 GPa or less, and

a standard deviation σ of the nanohardness of the sheet plane at the position of 1/2 of the depth of the soft surface layer from the surface of the high-strength steel sheet in the sheet thickness direction is 2.2 GPa or less.
The high-strength steel sheet according to any one of claims 1 to 4, comprising a metal coated or plated layer on a surface of the high-strength steel sheet on one side or both sides, the metal coated or plated layer containing one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi in a total amount of more than 50 mass%.
The high-strength steel sheet according to any one of claims 1 to 5, comprising a metal coated or plated layer in an outermost layer of the high-strength steel sheet on one side or both sides, the metal coated or plated layer containing at least one of zinc or aluminum in a total amount of 50 mass% or more.
A member produced using the high-strength steel sheet according to any one of claims 1 to 6.
An automotive framework structural component or automotive reinforcing component comprising the member according to claim 7.
A production method for a high-strength steel sheet, the production method comprising:
a hot rolling process of subjecting a steel slab having the chemical composition according to claim 1 or 3 to rough rolling under a condition that an average strain rate is 1 × 10^-4/s or more and 1 × 10^-1/s or less and a total rolling reduction ratio is 50 % or more, thereafter to finish rolling, and thereafter to coiling treatment to obtain a hot-rolled sheet;

thereafter a pickling and cold rolling process of subjecting the hot-rolled sheet to pickling and cold rolling to obtain a cold-rolled sheet;

thereafter a first heating process of subjecting the cold-rolled sheet to first heating under a condition that a heating temperature is 750 °C or more;

thereafter a first cooling process of cooling the cold-rolled sheet under a condition that a first cooling rate in a temperature range of T₂ or more and 750 °C or less is 2.0 °C/s or more;

thereafter an in-furnace retention process of retaining the cold-rolled sheet under a condition that a retention temperature T₂ is 350 °C or more and 550 °C or less and a retention time t in seconds satisfies that F defined in the following Formula 1 is 0.20 or more and 0.90 or less, $F = 1 - \exp (- {kt}^{n})$

where t is the retention time in seconds, and k and n are constants determined from an expansion curve of a Formaster test in which a test piece obtained by subjecting the slab to processes up to end of the first cooling process is held at the retention temperature T₂ of 350 °C or more and 550 °C or less;

thereafter a second cooling process of cooling the cold-rolled sheet to Ms - 20 °C or less under a condition that a second average cooling rate in a temperature range of Ms - 20 °C or more and Ms or less is 5 °C/s or more; and

thereafter a second heating process of treating the cold-rolled sheet under a condition that a temperature X in °C and a holding time Y in seconds satisfy the following Formula 2, $7000 \leq (273 + X) (20 + \log (Y / 3600)) \leq 13000 .$
The production method for a high-strength steel sheet according to claim 9, wherein in the second heating process, the temperature X in °C satisfies the following Formula 3, $100 \leq X \leq 400 .$
The production method for a high-strength steel sheet according to claim 9 or 10, wherein the first heating process is performed in an atmosphere with a dew point of -30 °C or more.
The production method for a high-strength steel sheet according to any one of claims 9 to 11, comprising a process of applying, to one side or both sides of the steel sheet after the cold rolling process and before the annealing process, metal coating or plating containing one or more selected from Cr, Mn, Fe, Co, Ni, Cu, Ga, Ge, As, Ru, Rh, Pd, Ag, Cd, In, Sn, Sb, Os, Ir, Rt, Au, Hg, Ti, Pb, and Bi in a total amount of more than 50 mass%.
The production method for a high-strength steel sheet according to any one of claims 9 to 12, comprising a process of applying, to the steel sheet during the process from the first heating to the second heating , metal coating or plating containing at least one of zinc or aluminum in a total amount of 50 mass% or more.
A production method for a member, the production method comprising subjecting the high-strength steel sheet according to any one of claims 1 to 6 to at least one of forming or joining to obtain the member.