EP4660346A1

EP4660346A1 - Steel sheet, member, and production methods for same

Info

Publication number: EP4660346A1
Application number: EP24779539.6A
Authority: EP
Inventors: Kotomi NOGUCHI; Yoichiro MATSUI; Tadachika Chiba; Hideyuki Kimura
Original assignee: JFE Steel Corp
Current assignee: JFE Steel Corp
Priority date: 2023-03-31
Filing date: 2024-03-14
Publication date: 2025-12-10
Also published as: JPWO2024203410A1; KR20250150125A; JP7635895B1; MX2025011270A; CN120858193A; WO2024203410A1

Abstract

Provided is a steel sheet and a member having a tensile strength (TS) of 980 MPa or more, high ductility and stretch flangeability, and excellent axial crash properties, and methods for producing the same.

A steel sheet has a chemical composition with specified C, Si, Mn, P, S, sol. Al, and N contents and has a steel microstructure containing, in area fractions, ferrite: 15% or less, tempered martensite: 30% or more and 80% or less, bainite: 10% or more and 40% or less, retained austenite: 5% or more, fresh martensite: 10% or less. The average C concentration in the retained austenite is 0.60 mass% or more. In a region within 100 µm from a steel sheet surface in a thickness direction, a content of Fe element present as carbides in the tempered martensite is 0.20 mass% or less on average.

Description

Technical Field

The present invention relates to a high-strength steel sheet that is applied to various usages including automobiles and home electric appliances and that has excellent formability suitable in particular for energy-absorbing members, and to a member and methods for producing the same.

Background Art

As the needs for weight reduction of automotive bodies have grown in recent years, high-strength steel sheets that have a tensile strength (TS) of 980 MPa or more have been increasingly applied to frame parts and seat parts of automobiles. However, when such high-strength steel sheets are formed by cold pressing, press cracking tends to occur due to degradation in ductility and stretch flangeability. Thus, these high-strength steel sheets are required to have formability superior to what has been typically achieved.
Furthermore, from the viewpoint of safety of passengers, deformation of the cabin and the vicinity thereof upon collision must be suppressed, and thus energy-absorbing members such as side members are required to absorb collision energy by becoming deformed upon collision. However, due to degradation in axial crash properties, high-strength steel sheets having a tensile strength of 980 MPa or more are prone to cause member fracture upon collision, starting from the site subjected to a primary process by forming, and have an issue of inability to stably exhibit the collision energy-absorbing function.
Thus, in order for high-strength steel sheets having TS of 980 MPa or more to be applicable to energy-absorbing members of automobiles, not only ductility and stretch flangeability but also excellent axial crash properties are required.
Under such circumstances, for example, Patent Literature 1 discloses that a steel sheet having a tensile strength of 980 MPa or more and excellent elongation and flangeability is obtained, the steel sheet containing, in mass%, C: 0.15% or more and 0.30% or less, P: 0.040% or less, S: 0.0100% or less, N: 0.0100% or less, O: 0.0060% or less, either or both of Si and Al: 0.70% or more and 2.50% or less in total, either or both of Mn and Cr: 1.50% or more and 3.50% or less in total, Mo: 0% or more and 1.00% or less, Ni: 0% or more and 1.00% or less, Cu: 0% or more and 1.00% or less, Nb: 0% or more and 0.30% or less, Ti: 0% or more and 0.30% or less, V: 0% or more and 0.30% or less, B: 0% or more and 0.0050% or less, Ca: 0% or more and 0.0400% or less, Mg: 0% or more and 0.0400% or less, and REM: 0% or more and 0.0400% or less, with the balance being Fe and impurities, in which the area fractions relative to the entire microstructure are as follows: either or both of ferrite and granular bainite: 10% or more and 50% or less in total, either or both of upper bainite and lower bainite: 10% or more and 50% or less in total, tempered martensite: more than 0% and 30% or less, retained austenite: 5% or more, and at least one of pearlite, cementite, and martensite: 0% or more and 10% or less in total and in which the area fraction of the ferrite relative to the total area fraction of the ferrite and the granular bainite is 25% or less.
Patent Literature 2 discloses that a high-strength hot-dip galvanized steel sheet having a tensile strength of 980 MPa or more and excellent fracture resistance upon collision is obtained, the steel sheet having a steel composition containing, in mass%, C: 0.07 to 0.20%, Si: 0.1 to 2.0%, Mn: 2.0 to 3.5%, P: 0.05% or less, S: 0.05% or less, and Sol. Al: 0.005 to 0.1% with the balance being Fe and incidental impurities, and having a steel microstructure in which area fractions are as follows: ferrite: 60% or less, tempered martensite: 40% or more, and fresh martensite: 10% or less, in which the void number-density in a bent portion in a VDA bending test is adjusted to 1500 voids/mm² or less.
Patent Literature 3 discloses that a high-strength cold rolled steel sheet having excellent strength, ductility, and flangeability is obtained, the steel sheet containing, in mass%, C: 0.10 to 0.40%, Mn: 0.5 to 4.0%, Si: 0.005 to 2.5%, Al: 0.005 to 2.5%, Cr: 0 to 1.0%, and the balance being iron and incidental impurities with P limited to 0.05% or less, S limited to 0.02%, and N limited to 0.006% or less, and having a steel microstructure containing, in area fractions, 2 to 30% of retained austenite and 20% or less of martensite, in which the average grain size of cementite is 0.01 µm or more and 1 µm or less, and the aforementioned cementite contains 30% or more and 100% or less of cementite having an aspect ratio of 1 or more and 3 or less.

Citation List

Patent Literature

PTL 1: Japanese Patent No. 6338038
PTL 2: Japanese Patent No. 6795122
PTL 3: Japanese Patent No. 4903915

Summary of Invention

Technical Problem

However, according to the technologies disclosed in Patent Literature 1 and Patent Literature 3, although a high-strength steel sheet having excellent ductility and stretch flangeability is provided, the axial crash properties are not considered.
Moreover, Patent Literature 2 discloses a high-strength steel sheet of 980 MPa or more having excellent fracture resistance upon collision; however, formability is not considered, and the technology is limited to a steel sheet having a hot-dip galvanized layer.
The present invention has been made to address the aforementioned issues, and an object thereof is to provide a steel sheet and a member having a tensile strength (TS) of 980 MPa or more, high ductility and stretch flangeability, and excellent axial crash properties, and methods for producing the same.
In the present invention, the tensile strength is measured by a tensile test according to JIS Z 2241 (2011).
In the present invention, high ductility means that the total elongation (T-El) as measured by a tensile test according to JIS Z 2241 (2011) is as follows: (A) in a TS range of 980 MPa or more and less than 1180 MPa, T-El is 14.0% or more; (B) in a TS range of 1180 MPa or more and less than 1320 MPa, T-El is 12.0% or more, and (C) in a TS range of 1320 MPa or more, T-El is 10.0% or more.
Furthermore, in the present invention, high stretch flangeability means that when a 100 mm × 100 mm square sample is punched with a punching tool having a punch diameter of 10 mm and a die diameter of 10.3 mm (13% clearance) and then the hole is expanded with a conical punch with 60 degree apex angle until cracks that penetrate through the thickness occur while arranging the burr, which has been formed by making the hole by punching, to face outward, the hole expansion ratio λ (%) = {(d - d₀)/d₀} × 100 is 45% or more where d₀ represents the initial hole diameter (mm) and d represents the hole diameter (mm) at the time the cracks occur.
In the present invention, excellent axial crash properties refer to a steel sheet having a VDA bend angle α of 80° or more in VDA bending.
Here, a VDA bending test is a bending test (Verband der Automobilindustrie: VDA bending test) according to VDA standards (VDA238-100) standardized by the German Association of the Automotive Industry, and is a 3-point bending test characterized by an extremely narrow roller spacing and a sharp punch. The VDA bending test uses a 60 mm × 60 mm square test specimen and involves supporting the test specimen by rollers having a roller diameter D of 30 mm and a roller spacing L of (thickness a₀ × 2) + 0.5 mm such that the bending ridge direction is parallel to the rolling direction, and then pushing a punch having a tip r of 0.4 mm into the test specimen from above at a stroke rate of 20 mm/min. The VDA bend angle α is an angle (° (more specifically, the unit is "°/mm" but hereinafter referred to as "°")) calculated by using equations (1) to (5) from the stroke S (mm) under the maximum load in the aforementioned bending test, and can be used as an indicator of the axial crash properties. Here, the stroke S (mm) under the maximum load is the length (mm) by which the punch has moved from the start of testing to the time point when the maximum load is obtained. $c (mm) = D / 2 + r + a_{0}$ $p (mm) = D / 2 + L / 2$
[Math. 1] $W (mm) = \sqrt{p^{2} + {(S - c)}^{2} - c^{2}}$
[Math. 2] $\sin \frac{αc}{2} = \frac{p \times c + W \times (S - c)}{p^{2} + {(S - c)}^{2}}$
[Math. 3] $\cos \frac{αc}{2} = \frac{W \times p - c \times (S - c)}{p^{2} + {(S - c)}^{2}}$

Solution to Problem

The inventors of the present invention have conducted extensive studies on a steel sheet having high strength, excellent ductility and stretch flangeability, and excellent axial crash properties, and have arrived at the following conclusions.

(1) A steel slab having a specified chemical composition is hot-rolled and cold-rolled, and then, from the viewpoints of strength, stretch flangeability, and axial crash properties, the resulting sheet is subjected to an annealing treatment in an austenite single-phase range involving an annealing temperature of (A_c3 temperature (°C) - 15°C) or higher and 900°C or lower and cooled from the annealing temperature at a specified cooling rate to control the ferrite area fraction to 15% or less.
(2) During the cooling performed after the aforementioned annealing treatment, the cooling stop temperature is set to 200°C or higher and (martensite transformation start temperature Ms (°C) - 30°C) or lower so as to partially substitute austenite with martensite and/or bainite. Martensite and bainite formed as such cause carbon to concentrate into the surrounding non-transformed austenite during the overaging treatment step after stopping the cooling and contribute to securing a specified C concentration in the retained austenite and improving the ductility.
In addition, it is possible to perform a retention treatment during cooling, and by retaining the temperature range from 500°C to a martensite transformation start temperature Ms or higher and 320°C or higher for 10 s or longer and 60 s or shorter, a larger amount of bainite can be obtained and the ductility can be improved further.
(3) In the aforementioned overaging treatment step after stopping the cooling, heating is performed to a heating temperature of 300°C or higher and 450°C or lower. Next, cooling is performed to a temperature 30°C to 150°C lower than the aforementioned heating temperature at an average cooling rate of 0.5°C/s or more, and further cooling is performed to a temperature of 150°C or higher and 350°C or lower at an average cooling rate of 0.1°C/s or more. Next, holding is performed in the temperature range from 150°C or higher and 350°C or lower for 20 s or longer and 1000 s or shorter. In this manner, precipitation of carbides in the tempered martensite present in the steel sheet surface layer and serving as the starting point of the crack propagation in the VDA bending can be minimized, and the axial crash properties can be improved. In addition, due to the aforementioned heat treatment step (the step of heating to 300 to 450°C after stopping the cooling and then finally holding the temperature range 150°C or higher and 350°C or lower for 20 s or longer and 1000 s or shorter), carbon partitioning from the surrounding microstructure to the non-transformed austenite region is accelerated, and thus the carbon concentration in the retained austenite can be further increased and the ductility can be improved.

As such, particularly by appropriately controlling the temperature history in the cooling process after the annealing and the overaging treatment process, decreasing the carbides in the tempered martensite and increasing the carbon concentration in the retained austenite can be simultaneously achieved. As a result, a steel sheet that has a tensile strength TS of 980 MPa or more and excellent ductility, stretch flangeability, and axial crash properties can be produced.
The present invention has been made based on the aforementioned findings and specifically provides the following.

[1] A steel sheet including:
- a chemical composition containing, in mass%,
  - C: 0.10% or more and 0.30% or less,
  - Si: 0.5% or more and 2.0% or less,
  - Mn: 1.5% or more and 3.0% or less,
  - P: 0.10% or less,
  - S: 0.020% or less,
  - sol. Al: 1.00% or less,
  - N: 0.015% or less, and
  - the balance being Fe and incidental impurities; and
- a steel microstructure containing, in area fractions, ferrite: 15% or less,
  - tempered martensite: 30% or more and 80% or less,
  - bainite: 10% or more and 40% or less,
  - retained austenite: 5% or more, and
  - fresh martensite: 10% or less,
- wherein:
  - an average C concentration in the retained austenite is 0.60 mass% or more, and
  - in a region within 100 µm from a steel sheet surface in a thickness direction, a content of Fe element present as carbides in the tempered martensite is 0.20 mass% or less on average.
[2] The steel sheet described in [1], wherein the chemical composition further contains, in mass%, at least one selected from
- Ti: 0.100% or less,
- Nb: 0.100% or less,
- V: 0.100% or less,
- B: 0.0050% or less,
- Cr: 1.000% or less,
- Cu: 1.000% or less,
- Sb: 0.200% or less,
- Sn: 0.200% or less,
- Ta: 0.100% or less,
- Ca: 0.0050% or less,
- Mg: 0.0050% or less, and
- REM: 0.0050% or less.
[3] The steel sheet described in [1] or [2], including a zinc coating layer on the steel sheet surface.
[4] A member including the steel sheet described in any one of [1] to [3].
[5] A method for producing a steel sheet, the method including:
- hot-rolling and cold-rolling a steel slab having the chemical composition described in [1] or [2] to obtain a cold rolled steel sheet;
- then subjecting the cold rolled steel sheet to an annealing treatment at an annealing temperature of (A_c3 temperature (°C) - 15°C) or higher and 900°C or lower;
- then cooling the cold rolled steel sheet in a range from the annealing temperature to a temperature T1 of 200°C or higher and (Ms temeperature (°C) - 30°C) or lower at an average cooling rate CR1 of 3°C/s or more and 100°C/s or less;
- heating the cold rolled steel sheet in a temperature range from the temperature T1 to a temperature T2 of 300°C or higher and 450°C or lower at an average heating rate of 2°C/s or more;
- cooling the cold rolled steel sheet in a temperature range from the temperature T2 to a temperature T3 of (T2 - 150°C) or higher and (T2 - 30°C) or lower at an average cooling rate CR2 of 0.5°C/s or more;
- cooling the cold rolled steel sheet in a temperature range from the temperature T3 to a temperature T4 of 150°C or higher and 350°C or lower and the temperature T3 or lower at an average cooling rate CR3 of 0.1°C/s or more;
- holding the cold rolled steel sheet in a temperature range from 150°C or higher and 350°C or lower and the temperature T4 or lower for 20 s or longer and 1000 s or shorter; and
- cooling the cold rolled steel sheet to a temperature of 50°C or lower at an average cooling rate CR4 of 1°C/s or more.
[6] A method for producing a steel sheet, the method including:
- hot-rolling and cold-rolling a steel slab having the chemical composition described in [1] or [2] to obtain a cold rolled steel sheet;
- then subjecting the cold rolled steel sheet to an annealing treatment at an annealing temperature of (A_c3 temperature (°C) - 15°C) or higher and 900°C or lower;
- then cooling the cold rolled steel sheet in a temperature range from the annealing temperature to 500°C at an average cooling rate CR5 of 5°C/s or more and 100°C/s or less;
- retaining the cold rolled steel sheet in a temperature range from 500°C to a retention stop temperature T5 of an Ms temperature (°C) or higher and 320°C or higher at an average cooling rate CR6 of 10°C/s or less for 10 s or longer and 60 s or shorter;
- cooling the cold rolled steel sheet in a temperature range from the retention stop temperature T5 to a temperature T1 of 200°C or higher and (Ms temperature (°C) - 30°C) or lower at an average cooling rate CR7 of 3°C/s or more and 100°C/s or less;
- heating the cold rolled steel sheet in a temperature range from the temperature T1 to a temperature T2 of 300°C or higher and 450°C or lower at an average heating rate of 2°C/s or more;
- cooling the cold rolled steel sheet in a temperature range from the temperature T2 to a temperature T3 of (T2 - 150°C) or higher and (T2 - 30°C) or lower at an average cooling rate CR2 of 0.5°C/s or more;
- cooling the cold rolled steel sheet in a temperature range from the temperature T3 to a temperature T4 of 150°C or higher and 350°C or lower and the temperature T3 or lower at an average cooling rate CR3 of 0.1°C/s or more;
- holding the cold rolled steel sheet in a temperature range from 150°C or higher and 350°C or lower and the temperature T4 or lower for 20 s or longer and 1000 s or shorter; and
- cooling the cold rolled steel sheet to a temperature of 50°C or lower at an average cooling rate CR4 of 1°C/s or more.
[7] The method for producing a steel sheet described in [5], further including hot-dip-galvanizing or hot-dip-galvannealing the steel sheet after the steel sheet is cooled at the average cooling rate CR1 of 3°C/s or more and 100°C/s or less.
[8] The method for producing a steel sheet described in [6], further including hot-dip-galvanizing or hot-dip-galvannealing the steel sheet after the steel sheet is cooled at the average cooling rate CR7 of 3°C/s or more and 100°C/s or less.
[9] The method for producing a steel sheet described in [5] or [6], further including electrogalvanizing the steel sheet after the steel sheet is cooled at the average cooling rate CR4 of 1°C/s or more.
[10] A method for producing a member, the method including a step of preparing a member by subjecting the steel sheet according to any one of [1] to [3] to either or both of a forming process and a joining process.

Advantageous Effects of Invention

According to the present invention, a steel sheet that has a tensile strength (TS) of 980 MPa or more and excellent ductility, stretch flangeability, and axial crash properties can be obtained.
The steel sheet of the present invention has excellent axial crash properties and thus is suitable for use in energy-absorbing members.
When the steel sheet obtained in the present invention is applied to automotive parts, automobile weight is reduced and the fuel efficiency is expected to improve.

Brief Description of Drawings

[Fig. 1] Fig. 1 is a microstructure image of one example of the steel sheet of the present invention.
[Fig. 2] Fig. 2 is a diagram illustrating a method for measuring a steel microstructure of the steel sheet of the present invention.
[Fig. 3] Fig. 3 is a diagram illustrating a method for producing the steel sheet of the present invention.
[Fig. 4] Fig. 4 is a diagram illustrating a method for calculating the VDA bend angle.

Description of Embodiments

The present invention will now be specifically described.
A steel sheet of the present invention includes a chemical composition containing, in mass%, C: 0.10% or more and 0.30% or less, Si: 0.5% or more and 2.0% or less, Mn: 1.5% or more and 3.0% or less, P: 0.10% or less, S: 0.020% or less, sol. Al: 1.00% or less, N: 0.015% or less, and the balance being Fe and incidental impurities; and a steel microstructure containing, in area fractions, ferrite: 15% or less, tempered martensite: 30% or more and 80% or less, bainite: 10% or more and 40% or less, retained austenite: 5% or more, and fresh martensite: 10% or less, wherein an average C concentration in the retained austenite is 0.60 mass% or more, and, in a region within 100 µm from a steel sheet surface in a thickness direction, a content of Fe element present as carbides in the tempered martensite is 0.20 mass% or less on average.
First, the chemical composition of the steel sheet of the present invention is described. In the description below, "%" indicating the unit of the content of a component element means "mass%" unless clearly indicated otherwise.

[C: 0.10% or more and 0.30% or less]

C is necessary for securing the tempered martensite amount and securing the retained austenite amount stable at room temperature, and, furthermore, stabilizes retained austenite by concentrating in the retained austenite and is thus a necessary element for improving the ductility. At a C content of less than 0.10%, these effects are not sufficiently exhibited, and it becomes difficult to obtain strength and workability of the steel sheet. The C content thus needs to be 0.10% or more, and is preferably 0.12% or more and more preferably 0.15% or more.
Meanwhile, at a C content exceeding 0.30%, the number of carbides in the tempered martensite increases, and cracks occur from the carbides as starting points during bending deformation; thus, the axial crash properties are degraded, and the ductility is also degraded due to an excessive increase in strength. The C content thus needs to be 0.30% or less, and is preferably 0.28% or less and more preferably 0.25% or less.

[Si: 0.5% or more and 2.0% or less]

Si suppresses formation of carbides in martensite and bainite, accelerates formation of retained austenite, and is an element useful for improving the stability of retained austenite. In order to obtain such effects, the Si content needs to be 0.5% or more, and is preferably 0.6% or more and more preferably 0.7% or more.
Meanwhile, at a Si content exceeding 2.0%, an excessive increase in strength degrades ductility and stretch flangeability, and, in addition, the rolling load may increase during hot rolling, liquid metal embrittlement cracking may occur when welded with a zinc-coated material, and the chemical convertibility may be degraded.
The Si content thus needs to be 2.0% or less, and is preferably 1.5% or less and more preferably 1.0% or less.

[Mn: 1.5% or more and 3.0% or less]

Mn is effective for securing the specified area fractions of tempered martensite and/or bainite and for securing the strength. Mn is also an important element for increasing the retained austenite area fraction and improving the ductility since Mn decreases the Ms temperature of the retained austenite and stabilizes the retained austenite. Thus, the Mn content needs to be 1.5% or more, and is preferably 1.8% or more and more preferably 2.0% or more.
Meanwhile, at a Mn content exceeding 3.0%, the strength increases excessively, bainite transformation delays extensively, and thus ductility is degraded. Moreover, the Mn segregation in the thickness direction becomes prominent, and there is a risk that the stability of the material quality would be degraded. The Mn content is thus 3.0% or less, and is preferably 2.8% or less and more preferably 2.5% or less.

[P: 0.10% or less]

P is an element that strengthen the steel, but a high P content degrades the spot weldability. Thus, the P content is 0.10% or less and is preferably 0.02% or less. It should be noted that P does not have to be contained but the P content is preferably 0.001% or more from the viewpoint of the production cost.

[S: 0.020% or less]

S is an element that has an effect of improving the descalability in hot rolling and an effect of suppressing nitriding during annealing, but also has extensive adverse effects on bendability, flangeability, and spot weldability. In order to mitigate these adverse effects, the S content is to be at least 0.020% or less. In the present invention, from the viewpoint of suppressing degradation of bendability and flangeability, the S content is preferably 0.0020% or less and more preferably less than 0.0010%. It should be noted that S does not have to be contained but the S content is preferably 0.0001% or more from the viewpoint of the production cost.
The S content is more preferably 0.0005% or more.

[sol. Al: 1.00% or less]

Al suppresses formation of carbides and is an element effective for accelerating formation of retained austenite. Moreover, Al is an element added as a deoxidizing agent in the steelmaking process. Thus, the lower limit of the sol. Al content is not particularly limited, but is preferably 0.005% or more and more preferably 0.01% or more for achieving stable deoxidization.
Meanwhile, at a sol. Al content exceeding 1.00%, inclusions in the steel sheet increase, and the ductility is degraded. Thus, the sol. Al content is 1.00% or less and is preferably 0.50% or less. The sol. Al content is more preferably 0.25% or less and even more preferably 0.10% or less.

[N: 0.015% or less]

N is an element that forms nitrides such as BN, AlN, and TiN in the steel and is an element that degrades the hot ductility of the steel and the surface quality. Moreover, a steel containing B has a negative effect of canceling the effects of B through formation of BN. At a N content exceeding 0.015%, the surface quality is extensively degraded. Thus, the N content is 0.015% or less and is preferably 0.010% or less. The N content is more preferably 0.005% or less and even more preferably 0.002% or less. It should be noted that N does not have to be contained but the N content is preferably 0.0001% or more and more preferably 0.001% or more from the viewpoint of the production cost.
The balance other than what is described above is Fe and incidental impurities. The steel sheet of the present invention preferably has a chemical composition that contains the aforementioned basic components and the balance being iron (Fe) and incidental impurities.
The chemical composition of the steel sheet of the present invention may contain the following optional elements as appropriate in addition to the aforementioned components.

[Ti: 0.100% or less, Nb: 0.100% or less, V: 0.100% or less]

Ti, Nb, and V form fine precipitates during hot rolling or annealing and thereby increase the strength. In order to obtain such an effect, the Ti content is preferably 0.010% or more, the Nb content is preferably 0.020% or more, and the V content is preferably 0.020% or more.
Meanwhile, when the Ti content, the Nb content, and the V content each exceed 0.100%, formability is degraded. Thus, when Ti, Nb, and V are to be added, each content is preferably 0.100% or less, more preferably 0.080% or less, and even more preferably 0.050% or less.

[B: 0.0050% or less]

B is an element that improves hardenability of the steel and has an advantage of facilitating formation of specified area fractions of tempered martensite and/or bainite. In order to obtain such an effect, the B content is preferably 0.0005% or more and even more preferably 0.0010% or more.
Meanwhile, at a B content exceeding 0.0050%, the effect thereof saturates, and the hot ductility is degraded extensively, resulting in surface defects. Thus, when B is to be added, the B content is preferably 0.0050% or less, more preferably 0.0040% or less, and even more preferably 0.0025% or less.

[Cr: 1.000% or less, Cu: 1.000% or less]

Cr and Cu not only serve as solid solution strengthening elements but also are elements that stabilize austenite and facilitate formation of a multi-phase microstructure during the cooling process in annealing. Thus, although Cr and Cu do not have to be contained, in order to obtain such effects, the Cr and Cu contents are each preferably 0.005% or more, more preferably 0.008% or more, and even more preferably 0.010% or more.
The Cr and Cu contents are each preferably 0.020% or more. The Cr and Cu contents are each more preferably 0.04% or more.
Meanwhile, at a Cr content and Cu content each exceeding 1.000%, formability of the steel sheet is degraded. Thus, when Cr and Cu are to be added, each content is preferably 1.000% or less, more preferably 0.800% or less, and even more preferably 0.400% or less. The Cr and Cu contents are each preferably 0.200% or less and more preferably 0.100% or less.

[Sb: 0.200% or less, Sn: 0.200% or less]

Sb and Sn are elements effective for suppressing decarburization in the region several tens of micrometers in the steel sheet surface layer formed by nitridation and oxidation of the steel sheet surface. When the nitridation and oxidation are suppressed, it becomes possible to prevent the amount of martensite in the steel sheet surface from decreasing and to secure the steel sheet strength and stability of the material quality. Thus, although Sb and Sn do not have to be added, but when Sb and Sn are to be added, the Sb and Sn contents are preferably each 0.002% or more, more preferably 0.004% or more, and even more preferably 0.006% or more.
Meanwhile, when these elements are each contained exceeding 0.200%, toughness is degraded. Thus, when Sb and Sn are to be added, the Sb content and the Sn content are each preferably 0.200% or less, more preferably 0.100% or less, and even more preferably 0.040% or less.

[Ta: 0.100% or less]

As with Ti and Nb, Ta forms alloy carbides and alloy carbonitrides and contributes to increasing the strength. In addition, it is considered that Ta has an effect of extensively suppressing coarsening of precipitates and an effect of stabilizing the contribution ratio to the steel sheet strength improvements by precipitation strengthening when Ta partly dissolves in the Nb carbides and Nb carbonitrides and forms complex precipitates such as (Nb,Ta) and (C,N). To obtain these effects, the Ta content is preferably 0.005% or more.
Meanwhile, excessive addition of Ta saturates the precipitate stabilizing effect and increases the alloying cost. Thus, when Ta is to be added, the Ta content is preferably 0.100% or less, more preferably 0.080% or less, and even more preferably 0.050% or less.

[Ca: 0.0050% or less, Mg: 0.0050% or less, REM: 0.0050% or less]

Ca, Mg, and REM are elements used in deoxidation, spheroidize the sulfides, and are elements effective for addressing adverse effects of the sulfides on the local ductility and the stretch flangeability. In order to obtain such effects, the Ca content is preferably 0.0001% or more, the Mg content is preferably 0.0001% or more, and the REM content is preferably 0.0001% or more.
Meanwhile, excessive addition of Ca, Mg, and REM exceeding 0.0050% induces the increase of inclusions etc., and causes defects and the like on the surface or in the inside. Thus, the Ca, Mg, and REM contents are each preferably 0.0050% or less, more preferably 0.0025% or less, and even more preferably 0.0010% or less.
The Ca content is more preferably 0.0008% or less.
The Mg content is more preferably 0.0008% or less.
The REM content is more preferably 0.0008% or less.
Here, REM in the present invention refers to scandium (Sc) having an atomic number 21, yttrium (Y) having an atomic number 39, and lanthanoids from lanthanum (La) having an atomic number 57 to lutetium (Lu) having an atomic number 71.
A REM content in the present invention is the total content of one or more of the elements selected from among the REMs described above. REM is not particularly limited but is preferably La and/or Ce.
When the aforementioned optional elements are contained in amounts lower than the preferable lower limit values, the optional elements contained in amount lower than the lower limit values do not impair the effects of the present invention. When the optional elements are contained in amounts lower than the lower limit values, the optional elements are considered incidental impurities.
Next, the steel microstructure (microstructure) of the steel sheet of the present invention is described.

[Ferrite area fraction: 15% or less]

Ferrite that is formed during annealing or the cooling process contributes to improving ductility; however, a difference in hardness occurs relative to the surrounding hard phases such as martensite, cracks propagate from the interface between ferrite and the hard phases during bending, and thus the stretch flangeability and the axial crash properties are degraded. Ferrite does not impair the effects of the present invention as long as the area fraction thereof is 15% or less, and thus the ferrite area fraction is 15% or less. The ferrite area fraction is preferably 10% or less and more preferably 5% or less. Furthermore, the ferrite area fraction may be 0%.

[Tempered martensite area fraction: 30% or more and 80% or less]

In order to obtain specified strength and stretch flangeability, the area fraction of tempered martensite is 30% or more. The area fraction of tempered martensite is preferably 40% or more and more preferably 50% or more.
Meanwhile, at a tempered martensite area fraction exceeding 80%, the ductility is degraded due to the excessive increase in strength, and thus the tempered martensite area fraction is 80% or less, preferably 75% or less, and more preferably 70% or less.

[Bainite: 10% or more and 40% or less]

Bainite in the present invention is a general term for upper bainite and lower bainite. Here, upper bainite and lower bainite are defined as follows. Upper bainite is composed of lath-like bainitic ferrite, and retained austenite and/or carbides present between laths of bainitic ferrite and is characterized by the absence of fine carbides regularly aligned within the lath-like bainitic ferrite. Meanwhile, as with upper bainite, lower bainite is composed of lath-like bainitic ferrite and retained austenite and/or carbides present between laths of bainitic ferrite, but lower bainite is characterized by the presence of fine carbides regularly aligned within the lath-like bainitic ferrite.
In other words, upper bainite and lower bainite are distinguished from each other by the absence or presence of fine carbides that are regularly aligned within bainitic ferrite.
Bainite that is formed during the cooling process after annealing and the retention process in the cooling process causes carbon to concentrate to the surrounding non-transformed austenite region and renders it possible to efficiently obtain more stable retained austenite. Thus, the bainite area fraction is 10% or more.
Meanwhile, since bainite is soft compared with tempered martensite, excessive formation of bainite decreases the strength and degrades the stretch flangeability due to the increase in the difference in hardness between bainite and tempered martensite. Thus, the bainite area fraction is 40% or less, preferably 35% or less, and more preferably 30% or less.
It should be noted that, compared with lower bainite, upper bainite causes concentration of a larger amount of carbon to the surrounding non-transformed austenite region, and thus is advantageous from the viewpoint of efficiently obtaining retained austenite. Thus, it is particularly preferable to contain 10% or more and more preferably 15% or more of upper bainite.
Meanwhile, since upper bainite is soft compared with lower bainite, excessive formation of upper bainite is disadvantageous from the viewpoint of improving strength and stretch flangeability. Thus, the upper bainite area fraction is preferably 30% or less and more preferably 25% or less.

[Retained austenite area fraction: 5% or more]

In order to secure high ductility as well as the desired strength, the retained austenite area fraction relative to the entire steel microstructure is 5% or more. The retained austenite area fraction is preferably 7% or more and more preferably 10% or more.
This amount of retained austenite includes retained austenite present between bainitic ferrite. The area fraction of the retained austenite present between bainitic ferrite can be measured in the same manner as area fractions of other retained austenite according to the measurement method described below.
Meanwhile, when the amount of retained austenite increases excessively, the strength may decrease, and the stretch flangeability and the delayed fracture resistance may be degraded. Thus, although the upper limit of the retained austenite area fraction is not specified, the retained austenite area fraction is preferably 20% or less and more preferably 15% or less.

[Fresh martensite area fraction: 10% or less]

The fresh martensite area fraction is 10% or less, preferably 8% or less, and more preferably 5% or less since fresh martensite may cause degradation of ductility and stretch flangeability. The fresh martensite area fraction may be 0%.

[Average C concentration in retained austenite: 0.60 mass% or more]

Increasing the C concentration in retained austenite stabilizes the retained austenite and makes it possible to improve ductility. Thus, in order to obtain sufficient ductility while securing the desired strength, the average C concentration in the retained austenite is 0.60 mass% or more and is preferably 0.80 mass% or more.
The upper limit is not particularly limited, but when the average C concentration in the retained austenite is excessively increased, the retained austenite does not undergo transformation during deformation, and the ductility improving effect may not be sufficiently obtained; thus the average C concentration in the retained austenite is preferably 2.0 mass% or less and more preferably 1.5 mass% or less.

[In a region within 100 µm from a steel sheet surface in the thickness direction, a content of Fe element present as carbides in tempered martensite: 0.20 mass% or less on average]

Carbides that precipitate in tempered martensite act as starting points of cracks during bending deformation and cause degradation of axial crash properties. In addition, precipitation of carbides is a phenomenon that competes with carbon partitioning, and since carbon is consumed by precipitation of the carbides, concentration of carbon into non-transformed austenite may be inhibited. Thus, the content of the Fe element present as carbides in the tempered martensite is preferably 0.20 mass% or less and more preferably 0.15 mass% or less on average. Although the lower limit is not particularly specified, in the present invention, the content of the Fe element is substantially 0.001 mass% or more.
The steel microstructure of the present invention contains, as described above, ferrite (including 0%), tempered martensite, bainite, retained austenite, and fresh martensite (including 0%). The steel microstructure of the present invention may include ferrite (including 0%), tempered martensite, bainite, retained austenite, and fresh martensite (including 0%).
Furthermore, the steel microstructure of the present invention may contain 5% or less of pearlite and cementite as phases other than those described above, and this does not impair the effects of the present invention.
Next, a method for measuring the steel microstructure is described.
The area fractions of ferrite, tempered martensite, and a massive microstructure (martensite austenite constituent (MA)) composed of fresh martensite and/or retained austenite described below are measured by cutting out a thickness section perpendicular to the steel sheet surface and parallel to the rolling direction, mirror-polishing the section, etching the mirror-polished section with a 3 vol% Nital, and observing ten 20 µm × 20 µm fields of view at a thickness 1/4 position by SEM at a magnification of 5000x. Fig. 1 illustrates one example of a SEM image of a steel microstructure of the steel sheet. As illustrated in Fig. 1, ferrite (in Fig. 1, see reference sign F) is a relatively equiaxed region where almost no carbides are found inside and which appears the darkest black by SEM. Tempered martensite (in Fig. 1, see reference sign TM) is a region where lath-like substructures and carbides are precipitated inside by SEM. MA (fresh martensite and/or retained austenite) (Fig. 1, see reference signs FM and RA) is a massive region where no substructure is seen inside and which appears white by SEM.
Furthermore, in the present invention, bainite is defined as the sum of the area fractions of upper bainite and lower bainite. Upper bainite is determined by observation by SEM as with ferrite etc., described above. A microstructure composed of bainitic ferrite, which is a region that does not involve precipitation of regularly aligned fine carbides and appears in the darkest black under SEM, and carbides or retained austenite that appears white under SEM is assumed to be upper bainite (in Fig. 1, see reference sign UB).
When it is difficult to distinguish ferrite from bainitic ferrite, a ferrite region with aspect ratio ≤ 2.0 is assumed to be ferrite and a region with aspect ratio > 2.0 is assumed to be bainitic ferrite. Fig. 2 is a diagram illustrating a method for measuring the steel microstructure of the steel sheet of the present invention. As illustrated in Fig. 2, the aspect ratio is determined as a/b irrespective of whether the shape is an elliptic shape illustrated in (1) or a polygonal shape illustrated in (2). Here, a is the length of the longest segment among the segments formed by any two points on the region perimeter of each phase, and b is the length of the segment that is perpendicular to the straight line forming a and that is the longest among the segments formed by any two points on the region perimeter.
In addition, as illustrated in (3) of Fig. 2, when determining the aspect ratio for the case in which grains are in contact with each other, measurement is performed without considering the grains as one grain but separately as two or more grains.
Specifically, when grains are in contact with each other, as illustrated by a broken line position in (3) of Fig. 2, division is made at a position where the area fraction of each grain is about the same, and then the size of each grain is measured.
In addition, under SEM observation, lower bainite (in Fig. 1, see reference sign LB) and tempered martensite both appear to have microstructures where fine white carbides are precipitated on a grey background, and thus are difficult to distinguish from each other. Thus, lower bainite and tempered martensite are distinguished by observing the variant forms of carbides by transmission electron microscopy (TEM). Here, the carbide form of lower bainite is a single variant in which precipitation occurs regularly in one direction inside the substructure, whereas the carbides of tempered martensite are a multivariant in which precipitation directions are random inside the substructure. The area fractions of lower bainite and tempered martensite having such characteristics can be determined by observing ten 1.5 µm square fields of view by TEM, calculating the area fractions of the constituent phases (lower bainite and tempered martensite) by using Adobe Photoshop described above from ten fields of view, and averaging the obtained values.
The area fraction of retained austenite and the average C concentration in retained austenite are determined by grinding and polishing the steel sheet in the thickness direction down to a thickness 1/4 position and performing X-ray diffraction intensity measurement thereon. The volume fraction of retained austenite is calculated by using a Mo tubular radiation source from the intensity ratios of the (200), (220), and (311) planes of austenite relative to the diffraction intensities of the (200) and (211) planes of ferrite. The average C concentration in retained austenite is determined by using a Co tubular radiation source, determining the lattice constant A of austenite from the peak angle of the (220) plane of austenite, and calculating the concentration from equation (6) below. Note that, in the equation below, Mn%, Si%, and Al% each indicate mass% of the content of each element (Mn, Si, and Al) in the steel sheet. Average C concentration in retained austenite (mass%) = A - {(0.3572 + 0.0012 × Mn% - 0.00157 × Si% + 0.0056 × Al%)}/0.033
The area fraction of fresh martensite (FM) is determined from equation (7) below, and is a value obtained by subtracting the area fraction of retained austenite (RA) from the area fraction of the aforementioned MA (fresh martensite and/or retained austenite).
In the present invention, the "area fraction" of RA (retained austenite) can be regarded as equivalent to the "volume fraction" of retained austenite determined by XRD measurement. $FM (%) = MA (%) - RA (%)$
In the region within 100 µm from the steel sheet surface in the thickness direction, the average content of the Fe element present as carbides in tempered martensite is determined by extracted residue analysis. A 10 mass% AA electrolyte (acetylacetone-1 mass% tetramethylammonium chloride-methanol) is used as the electrolyte, and electrolysis is performed under conditions of current density of 20 mA/cm² and an electrolysis time of 30 min to dissolve a region within 100 µm from the steel sheet surface. Subsequently, the specimen is immersed in methanol, the residue separated by ultrasonic agitation (precipitates that remain undissolved by electrolysis) is filtered out by using a filter having a 0.2 µm pore diameter, and the Fe element amount in the collected residue is determined by ICP atomic emission spectroscopy and is converted into a mass% in the steel sheet. Note that the collected residue is solely composed of precipitates that are present in the region within 100 µm from the steel sheet surface, and, here, the number of carbides outside the tempered martensite is significantly small (bainite, ferrite, etc.) compared with the number of carbides in the tempered martensite. Moreover, it is assumed that the Fe element does not exist as iron-based compounds other than carbon-based compounds. Thus, by the aforementioned measurement, the average content of the Fe element present as carbides in tempered martensite in the region within 100 µm from the steel sheet surface in the thickness direction is obtained.
The steel sheet of the present invention described heretofore may be a steel sheet that has a zinc coating layer on a steel sheet surface (one surface or both surfaces). The coating layer may be a hot-dip coating layer or an electroplating layer.
The thickness of the steel sheet of the present invention is preferably 0.5 mm or more. The thickness is preferably 2.0 mm or less.
Next, a method for producing the steel sheet of the present invention is described.
In the description below, the case where a retention treatment under specified conditions is not performed as the treatment that follows the annealing treatment is described as a first embodiment, and the case where the retention treatment under specified conditions is performed as the treatment that follows the annealing treatment is described as a second embodiment.
Here, temperatures specified in each step of the present invention are surface temperatures of slabs (steel slabs) or steel sheets.
Fig. 3 is a diagram illustrating the method for producing the steel sheet of the present invention, in particular, changes with time in surface temperature of a slab (steel slab) or a steel sheet. Details of the individual steps including the changes with time in temperature will now be described. Fig. 3(a) illustrates the changes with time in surface temperature of the slab (steel slab) or the steel sheet in the method for producing the steel sheet according to the first embodiment (the case in which the retention treatment is not performed). Meanwhile, Fig. 3(b) illustrates the changes with time in surface temperature of the slab (steel slab) or the steel sheet in the method for producing the steel sheet according to the second embodiment (the case in which the retention treatment is performed).

As illustrated in Fig. 3(a), according to a method for producing a steel sheet according to the first embodiment of the present invention, a steel slab having the aforementioned chemical composition is hot-rolled and cold-rolled, and the obtained cold rolled steel sheet is subjected to an annealing treatment at an annealing temperature of (A_c3 temperature (°C) - 15°C) or higher and 900°C or lower, is cooled at an average cooling rate CR1 of 3°C/s or more and 100°C/s or less in a range from the annealing temperature to a temperature T1 of 200°C or higher and (Ms temperature (°C) - 30°C) or lower, is heated at an average heating rate of 2°C/s or more in a temperature range from the temperature T1 to a temperature T2 of 300°C or higher and 450°C or lower, is cooled at an average cooling rate CR2 of 0.5°C/s or more in a temperature range from the temperature T2 to a temperature T3 of (T2 - 150°C) or higher and (T2 - 30°C) or lower, is cooled at an average cooling rate CR3 of 0.1°C/s or more in a temperature range from the temperature T3 to a temperature T4 of 150°C or higher and 350°C or lower and the temperature T3 or lower, is held in a temperature range 150°C or higher and 350°C or lower and the temperature T4 or lower for 20 s or longer and 1000 s or shorter, and is cooled at an average cooling rate CR4 of 1°C/s or more to a temperature of 50°C or lower.

(Hot rolling step)

Hot rolling may be performed according to a standard method; for example, a steel slab having the aforementioned specified chemical composition may be heated to a slab heating temperature of 1100°C or higher, hot-rolled by setting the soaking time to 20 min or longer and a finish rolling delivery temperature to the Ar₃ transformation temperature or higher, and coiled at a coiling temperature of 400°C or higher.
The slab heating temperature may be 1300°C or lower. The soaking time may be 300 min or shorter. The finish rolling delivery temperature may be equal to or lower than Ar₃ transformation temperature + 200°C. The coiling temperature may be 720°C or lower.
The coiling temperature is preferably controlled from the viewpoints of suppressing the thickness variation and stably securing high strength. Specifically, the coiling temperature is preferably 430°C or higher. The coiling temperature is preferably 530°C or lower.

(Cold rolling step)

Cold rolling may be performed at a rolling reduction (cumulative rolling reduction) of 30% or more. The reduction may be 85% or less. The rolling reduction is preferably controlled from the viewpoints of stably obtaining high strength and decreasing anisotropy. Specifically, the rolling reduction is preferably 35% or more. Note that when the rolling load is high, it is possible to perform an annealing treatment for softening in a continuous annealing line (CAL) or a box annealing furnace (BAF) at 450°C or higher and 730°C or lower.

(Annealing step)

The steel slab having the specified chemical composition is hot-rolled and cold-rolled and then annealed under the conditions specified below. The annealing facility is not particularly limited; however, from the viewpoints of productivity and achieving the desired heating rates and cooling rates, a continuous annealing line (CAL) or a continuous galvanizing line (CGL) is preferably used.

[Annealing treatment step: annealing temperature: (A_c3 temperature (°C) - 15°C) or higher and 900°C or lower]

In order to secure specified area fractions of tempered martensite and/or bainite and a specified area fraction of retained austenite, the annealing temperature is A_c3 temperature (°C) - 15°C or higher and 900°C or lower. In order to adjust the ferrite area fraction to 15% or less, the annealing temperature is preferably adjusted so that annealing occurs in the austenite single-phase range. The annealing temperature is preferably the A_c3 temperature or higher.
The annealing temperature is more preferably equal to or higher than A_c3 temperature (°C) + 15°C.
Meanwhile, at an annealing temperature exceeding 900°C, the austenite grain size increases excessively, the diffusion distance of C atoms necessary for obtaining retained austenite having a desired carbon concentration increases, and thus the ductility is degraded. Thus, the annealing temperature is 900°C or lower. The annealing temperature is preferably 880°C or lower.
Here, the A_c3 temperature can be determined by measuring the change in volume observed when a columnar test specimen (diameter of 3 mm × height of 10 mm) is heated from room temperature (25°C) to the austenite single-phase range by using a Formastor tester.

[First cooling treatment step: cooling at an average cooling rate CR1 of 3°C/s or more and 100°C/s or less in a range from the annealing temperature to a temperature T1 of 200°C or higher and (martensite transformation start temperature Ms (°C) - 30°C) or lower]

After the aforementioned annealing treatment, cooling must be rapidly performed in order to obtain the specified area fractions of tempered martensite and/or bainite. When the average cooling rate CR1 in the temperature range from the annealing temperature to a temperature T1 (cooling stop temperature T1) of 200°C or higher and (martensite transformation start temperature Ms (°C) - 30°C) or lower is less than 3°C/s, ferrite is formed during the cooling process, and this degrades strength, stretch flangeability, and axial crash properties. Thus, from the viewpoint of suppressing formation of ferrite, the average cooling rate CR1 in the temperature range from the annealing temperature to a cooling stop temperature T1 of 200°C or higher and (Ms temperature (°C) - 30°C) or lower is 3°C/s or more. The average cooling rate CR1 is preferably 5°C/s or more and more preferably 8°C/s or more.
Meanwhile, at an excessively large average cooling rate CR1, the sheet shape is degraded; thus, the average cooling rate CR1 is 100°C/s or less. The average cooling rate CR1 is preferably 50°C/s or less.
Furthermore, in order to obtain the specified retained austenite amount, the temperature T1 (cooling stop temperature T1) is 200°C or higher. The temperature T1 (cooling stop temperature T1) is preferably 220°C or higher and more preferably 240°C or higher. At a cooling stop temperature T1 exceeding (Ms temperature (°C) - 30°C), a large amount of massive non-transformed austenite remains, the amount of fresh martensite in the final cooling increases, and the stretch flangeability is degraded. Thus, the cooling stop temperature T1 is (Ms temperature (°C) - 30°C) or lower. The cooling stop temperature T1 is preferably (Ms temperature (°C) - 35°C) or lower.
Here, the martensite transformation start temperature Ms (°C) can be determined with a Formastor tester by measuring the change in volume of a columnar test specimen, which is similar to the one used in measuring the A_c3 temperature, observed when the test specimen is held at a specified annealing temperature and then rapidly cooled with helium gas.
Here, the average cooling rate CR1 is "(annealing temperature (cooling start temperature) (°C) - (Ms temperature (°C) - 30°C) (cooling stop temperature (°C)))/(cooling time (s) from annealing temperature to (Ms temperature (°C) - 30°C))".

[Heat treatment step: heating at an average heating rate of 2°C/s or more from the temperature T1 to a temperature T2 of 300°C or higher and 450°C or lower]

When the temperature range from the aforementioned temperature T1 (cooling stop temperature T1: 200°C or higher and (martensite transformation start temperature Ms (°C) - 30°C) or lower) to a temperature T2 of 300°C or higher and 450°C or lower is heated in a short period of time, carbide precipitation can be suppressed, and, simultaneously, carbon from the surrounding microstructure can concentrate into the non-transformed austenite, and thus high ductility can be obtained. When the temperature T2 is lower than 300°C, carbon does not sufficiently concentrate into non-transformed austenite, tempering of martensite does not progress sufficiently, and tempered martensite exhibits an excessively increased strength; thus, such a temperature T2 is disadvantageous from the viewpoint of improving ductility and stretch flangeability.
Meanwhile, when the temperature T2 after heating exceeds 450°C, non-transformed austenite decomposes, and massive fresh martensite increases; thus, the ductility and stretch flangeability are degraded, carbides precipitate excessively, and the axial crash properties are degraded. Thus, the temperature T2 is 300°C or higher and 450°C or lower. The temperature T2 is preferably 330°C or higher and more preferably 350°C or higher. The temperature T2 is preferably 420°C or lower and more preferably 400°C or lower.
Furthermore, when the average heating rate from the temperature T1 to a temperature T2 of 300°C or higher and 450°C or lower is slow, the axial crash properties are degraded by the increase in the number of carbide precipitates, the concentrated carbon amount in retained austenite decreases with the increase in carbides, and thus the ductility is degraded. Thus, the average heating rate in the temperature range from the temperature T1 (cooling stop temperature T1) to a temperature T2 of 300°C or higher and 450°C or lower is 2°C/s or more. From the viewpoint of suppressing carbide precipitation, the average heating rate is preferably 5°C/s or more and more preferably 10°C/s or more. Although the upper limit of the average heating rate is not particularly limited, the average heating rate is preferably 50°C/s or less and more preferably 30°C/s or less.
Here, the average heating rate is "(temperature T2 (heating stop temperature) (°C) - temperature T1 (heating start temperature) (°C))/(heating time (s) from temperature T1 to temperature T2)".

[Second cooling treatment step: cooling at an average cooling rate CR2 of 0.5°C/s or more in the temperature range from the temperature T2 to a temperature T3 of (T2 - 150°C) or higher and (T2 - 30°C) or lower]
[Third cooling treatment step: cooling at an average cooling rate CR3 of 0.1°C/s or more in the temperature range from the temperature T3 to a temperature T4 of 150°C or higher and 350°C or lower and the temperature T3 or lower]

In the overaging treatment step after the stop of cooling (from the aforementioned heat treatment step to the holding treatment step after the third cooling treatment step), fresh martensite turns into tempered martensite, and, with this transformation, precipitation of carbides and carbon partitioning to non-transformed austenite occur. The carbide precipitation rate tends to be slow as the heating temperature after the stop of cooling decreases; thus, it becomes possible to suppress carbide precipitation and improve the axial crash properties by rapid cooling at a particular cooling rate or more to a temperature range in which suppression of carbide precipitation is possible. Meanwhile, as the heating temperature after the stop of cooling decreases, concentration of carbon into the non-transformed austenite and tempering of fresh martensite are inhibited, and thus the strength may excessively increase and the ductility and stretch flangeability may be degraded. Thus, in order to simultaneously achieve excellent axial crash properties, ductility, and stretch flangeability by suppressing carbide precipitation by cooling to a specified temperature range and by accelerating concentration of carbon in non-transformed austenite and tempering of fresh martensite, a second cooling treatment step that involves specified conditions is performed between the heat treatment step and the third cooling treatment step.
Specifically, in order to obtain the aforementioned effects, the temperature T3 (cooling stop temperature T3) in the second cooling treatment step is (T2 - 150°C) or higher and (T2 - 30°C) or lower. The temperature T3 is preferably (T2 - 120°C) or higher and more preferably (T2 - 100°C) or higher. The temperature T3 is preferably (T2 - 50°C) or lower and more preferably (T2 - 70°C) or lower.
The average cooling rate CR2 is 0.5°C/s or more, preferably 0.6°C/s or more and more preferably 0.7°C/s or more. Although the upper limit is not particularly specified, a higher average cooling rate delays carbon partitioning to retained austenite and tempering of fresh martensite, and this tends to degrade ductility and stretch flangeability; thus, the average cooling rate CR2 is preferably 5°C/s or less.
Here, the average cooling rate CR2 is "(temperature T2 (°C) (cooling start temperature) - temperature T3 (°C) (cooling stop temperature))/(cooling time (s) from temperature T2 to temperature T3)".
After cooling to the temperature T3 as the second cooling treatment, in order to accelerate carbon partitioning simultaneously with suppression of carbide precipitation and to improve the carbon concentration in retained austenite and the retained austenite amount, further cooling is performed as a third cooling treatment. During this process, when the cooling temperature is decreased, the carbon partitioning rate decreases, and carbon may not sufficiently concentrate to non-transformed austenite.
Meanwhile, in response to the increase in cooling temperature and the decrease in average cooling rate, carbides precipitate during cooling and the axial crash properties are degraded. Thus, the temperature T4 is 150°C or higher and 350°C or lower. The temperature T4 is preferably 180°C or higher and more preferably 200°C or higher. In addition, the temperature T4 is preferably 300°C or lower and more preferably 250°C or lower.
The average cooling rate CR3 is 0.1°C/s or more, preferably 0.2°C/s or more, and more preferably 0.3°C/s or more. Although the upper limit is not particularly specified, from the viewpoint of sufficiently accelerating carbon partitioning to non-transformed austenite, the average cooling rate CR3 is preferably average cooling rate CR2 or less and 1°C/s or less.
Here, the average cooling rate CR3 is "(temperature T3 (°C) (cooling start temperature) - temperature T4 (°C) (cooling stop temperature))/(cooling time (s) from temperature T3 to temperature T4)".

[Holding treatment step: holding the temperature range from 150°C or higher and 350°C or lower and the temperature T4 or lower for 20 s or longer and 1000 s or shorter]

The holding temperature is the aforementioned temperature T4 or lower.
At a holding temperature of lower than 150°C, tempering of fresh martensite does not progress sufficiently, an excessive increase in strength occurs, and the concentrated carbon amount in the non-transformed austenite decreases; thus, such a holding temperature is disadvantageous from the viewpoint of improving ductility and stretch flangeability. Meanwhile, at a holding temperature exceeding 350°C, the precipitation number of carbides increases, the axial crash properties are degraded, concentration of carbon into non-transformed austenite is impaired by the carbide precipitation, retained austenite decreases, and the ductility is degraded.
Thus, the holding temperature is 150°C or higher and 350°C or lower and the temperature T4 or lower. The holding temperature is preferably 180°C or higher and more preferably 200°C or higher. The holding temperature is preferably 300°C or lower and more preferably 250°C or lower.
Extending the holding time increases the concentrated carbon amount in non-transformed austenite, can stabilize retained austenite, and can improve the ductility. Meanwhile, shortening the holding time tends to decrease the precipitation number of carbides and improves the axial crash properties. Thus, the holding time is 20 s or longer and 1000 s or shorter. The holding time is preferably 100 s or longer and more preferably 200 s or longer. The holding time is preferably 700 s or shorter and more preferably 500 s or shorter.
In addition, it is possible to form a zinc coating layer on a steel sheet surface by performing a hot-dip galvanizing treatment and even an alloying treatment after the first cooling treatment step (after cooling at an average cooling rate CR1 of 3°C/s or more and 100°C/s or less), that is, before, after, or during any one of the steps from the heat treatment step of heating from the temperature T1 to the temperature T2 to the holding treatment step of holding the temperature at 150°C or higher and 350°C or lower. In such a case, the steel sheet is preferably dipped in a zinc coating bath at 440°C or higher and 500°C or lower to perform a hot-dip galvanizing treatment, and then the coating weight is preferably adjusted by gas wiping or the like. The galvanization preferably uses a zinc coating bath having an Al content of 0.10 mass% or more and 0.22 mass% or less. Moreover, an alloying treatment of the zinc coating can be performed after the hot-dip galvanizing treatment. The alloying treatment on the zinc coating is preferably performed in the temperature range of 470°C or higher and 590°C or lower, and heating to this temperature range does not impair the effects of the present invention and thus may be performed.

[Fourth cooling treatment step: cooling to a temperature of 50°C or lower at an average cooling rate CR4 of 1°C/s or more]

Subsequently, cooling is performed at an average cooling rate CR4 of 1°C/s or more from the holding end temperature (150°C or higher and 350°C or lower and the temperature T4 or lower) in the aforementioned holding treatment step to a temperature of 50°C or lower from the viewpoint of preventing softening caused by excessive tempering and degradation of the ductility caused by carbide precipitation. From the viewpoint of adjusting the surface roughness and stabilizing the press formability such as flattening the sheet shape and from the viewpoint of increasing the yield strength, the steel sheet may be skin-pass rolled. The skin-pass elongation is preferably 0.1% or more and 0.5% or less. Furthermore, the sheet shape can be flattened by using a leveler.
The average cooling rate CR4 from the aforementioned holding end temperature to a temperature of 50°C or lower is 1°C/s or more and preferably 5°C/s or more. In the present invention, the upper limit of the average cooling rate CR4 is not particularly specified; however, from the viewpoint of production and the viewpoint of suppressing transformation from retained austenite to fresh martensite, the average cooling rate CR4 is preferably 10°C/s or less.
Here, the average cooling rate CR4 is "(150°C or higher and 350°C or lower and temperature T4 (°C) or lower (holding end temperature (cooling start temperature)) - temperature of 50°C or lower (cooling stop temperature))/(cooling time (s) from holding end temperature to cooling stop temperature)".
From the viewpoint of improving stretch flangeability, after the aforementioned heat treatment or skin-pass rolling, a low-temperature heat treatment can be performed at 100 to 300°C for 30 s to 10 days. By this treatment, hydrogen that has penetrated the steel sheet during tempering of the martensite formed during the final cooling or skin-pass rolling and annealing is removed from the steel sheet. The low-temperature heat treatment can decrease hydrogen to less than 0.1 mass ppm.
In addition, electroplating can be performed. That is, the steel sheet may be electrogalvanized after the aforementioned fourth cooling treatment step (after cooling at an average cooling rate CR4 of 1°C/s or more). From the viewpoint of decreasing hydrogen in the steel, the aforementioned low-temperature heat treatment is preferably performed after electroplating.

As illustrated in Fig. 3(b), according to a method for producing a steel sheet according to the second embodiment of the present invention, a steel slab having the aforementioned chemical composition is hot-rolled and cold-rolled, and the obtained cold rolled steel sheet is subjected to an annealing treatment at an annealing temperature of (A_c3 temperature (°C) - 15°C) or higher and 900°C or lower, is cooled at an average cooling rate CR5 of 5°C/s or more and 100°C/s or less in the temperature range from the annealing temperature to 500°C (pre-retention treatment first cooling treatment), is retained for 10 s or longer and 60 s or shorter at an average cooling rate CR6 of 10°C/s or less in the temperature range from 500°C to a retention stop temperature T5 of the Ms temperature (°C) or higher and 320°C or higher (retention treatment), is cooled at an average cooling rate CR7 of 3°C/s or more and 100°C/s or less in a temperature range from the retention stop temperature T5 to a temperature T1 of 200°C or higher and (Ms temperature (°C) - 30°C) or lower (post-retention treatment first cooling treatment), is heated at an average heating rate of 2°C/s or more in a temperature range from the temperature T1 to a temperature T2 of 300°C or higher and 450°C or lower, is cooled at an average cooling rate CR2 of 0.5°C/s or more in a temperature range from the temperature T2 to a temperature T3 of (T2 - 150°C) or higher and (T2 - 30°C) or lower, is cooled at an average cooling rate CR3 of 0.1°C/s or more in a temperature range from the temperature T3 to a temperature T4 of 150°C or higher and 350°C or lower and the temperature T3 or lower, is held in a temperature range of 150°C or higher and 350°C or lower and the temperature T4 or lower for 20 s or longer and 1000 s or shorter, and is cooled at an average cooling rate CR4 of 1°C/s or more to a temperature of 50°C or lower.
In the second embodiment, hot rolling, cold rolling, and the annealing treatment can be performed under the same conditions as those in the first embodiment.
Furthermore, in the second embodiment, the treatment in the first cooling treatment step of the first embodiment is replaced by the pre-retention treatment first cooling treatment, the retention treatment, and the post-retention treatment first cooling treatment.
Furthermore, in the second embodiment, the heat treatment, the second cooling treatment, the third cooling treatment, the holding treatment, and the fourth cooling treatment after the post-retention treatment first cooling treatment can be performed under the same conditions as those of the heat treatment, the second cooling treatment, the third cooling treatment, the holding treatment, and the fourth cooling treatment of the first embodiment, respectively. In addition, the hot-dip galvanizing treatment in the second embodiment can be performed under the same conditions as those of the first embodiment except that, whereas the hot-dip galvanizing treatment is performed after the first cooling treatment step (after cooling at an average cooling rate CR1 of 3°C/s or more and 100°C/s or less) in the first embodiment, the hot-dip galvanizing treatment is performed after the post-retention treatment first cooling treatment step (after cooling at an average cooling rate CR7 of 3°C/s or more and 100°C/s or less) in the second embodiment. Furthermore, in the second embodiment, other conditions including those of the low-temperature heat treatment after all of the heat treatments and the electroplating treatment can be the same as those of the first embodiment. In this embodiment, the pre-retention treatment first cooling treatment, the retention treatment, and the post-retention treatment first cooling treatment are mainly described below.

[Pre-retention treatment first cooling treatment step: cooling at an average cooling rate CR5 of 5°C/s or more and 100°C/s or less in a temperature range from the annealing temperature to 500°C]

[Retention treatment: retaining at an average cooling rate CR6 of 10°C/s or less for 10 s or longer and 60 s or shorter in a temperature range from 500°C to a retention stop temperature T5 of the Ms temperature (°C) or higher and 320°C or higher]

By performing the retention treatment during cooling, a larger area fraction of bainite is obtained, carbon concentrates into non-transformed austenite from this bainite, the retained austenite becomes stable, and thus a steel sheet having more superior ductility can be obtained. Thus, from the viewpoint of increasing ductility, the retention treatment is preferably performed during cooling.
When the retention treatment is performed during cooling, first, the temperature range from the annealing temperature to 500°C is cooled at an average cooling rate CR5 of 5°C/s or more and 100°C/s or less.
When the average cooling rate CR5 is less than 5°C/s, a large amount of ferrite is formed, and the strength and λ are degraded. The average cooling rate CRS is preferably 8°C/s or more.
Meanwhile, at an excessively high average cooling rate CR5, the sheet shape is degraded; thus, the average cooling rate CR5 is 100°C/s or less, preferably 50°C/s or less, and more preferably less than 30°C/s.
Here, the average cooling rate CR5 is "(annealing temperature (cooling start temperature) (°C) - 500°C (cooling stop temperature))/cooling time (s) from annealing temperature to 500°C".
Next, the temperature range from 500°C to a retention stop temperature T5 of the martensite transformation start temperature Ms (°C) or higher and 320°C or higher is retained at an average cooling rate CR6 of 10°C/s or less for 10 s or longer and 60 s or shorter; in this manner, bainite can be formed, and retained austenite having a high C concentration can be formed adjacent thereto. When the temperature range is lower than the Ms temperature (°C), martensite is formed preferentially, bainite transformation progresses excessively due to the swing back phenomenon, and the strength and the stretch flangeability are degraded. Meanwhile, when the aforementioned temperature range exceeds 500°C, the drive force of the bainite transformation decreases, and the bainite transformation amount decreases. In addition, when the aforementioned temperature range is lower than 320°C, lower bainite is mainly formed; however, from the viewpoint of improving ductility and VDA bending properties, upper bainite that does not accompany carbide precipitation is advantageous from the viewpoint of suppressing carbon-concentrated carbides in the non-transformed γ caused by carbide precipitation, and of degradation of the axial crash properties. Thus, the retention stop temperature T5 is 320°C or higher. Thus, the temperature range is the Ms temperature (°C) or higher and 320°C or higher and 500°C or lower. The aforementioned temperature range is preferably 350°C or higher and more preferably 380°C or higher. The aforementioned temperature range is preferably 480°C or lower and more preferably 460°C or lower.
The bainite transformation amount decreases when the average cooling rate CR6 exceeds 10°C/s. Thus, the average cooling rate CR6 is 10°C/s or less. Furthermore, when the retention time is shorter than 10 s, the desired amount of bainite is not obtained, and when the retention time is longer than 60 s, C concentrates into the massive non-transformed austenite from the bainite, and the amount of remaining massive microstructure increases. Thus, the retention time is 10 s or longer and 60 s or shorter. From the viewpoint of securing bainite and retained austenite, the retention time is preferably 20 s or longer. In addition, from the viewpoint of improving the stretch flangeability by decreasing the massive microstructure, the retention time is preferably 50 s or shorter.
Here, the average cooling rate CR6 is "(500°C (retention start temperature) (°C) - (retention stop temperature (°C) of Ms temperature (°C) or higher and 320°C or higher))/(retention time (s) from retention start temperature to retention stop temperature)".

[Post-retention treatment first cooling treatment step: cooling at an average cooling rate CR7 of 3°C/s or more and 100°C/s or less in a range from the retention stop temperature T5 to a temperature T1 of 200°C or higher and (martensite transformation start temperature Ms (°C) - 30°C) or lower]

After the aforementioned retention treatment, cooling must be rapidly performed in order to obtain the specified area fractions of tempered martensite and/or bainite. When the average cooling rate CR7 in the temperature range from the retention stop temperature T5 to a temperature T1 (cooling stop temperature T1) of 200°C or higher and (martensite transformation start temperature Ms (°C) - 30°C) or lower is less than 3°C/s, ferrite is formed during the cooling process, and this degrades the strength, stretch flangeability, and axial crash properties.
Thus, from the viewpoint of suppressing formation of ferrite, the average cooling rate CR7 in the temperature range from the retention stop temperature T5 to the cooling stop temperature T1 of 200°C or higher and (Ms temperature (°C) - 30°C) or lower is 3°C/s or more. The average cooling rate CR7 is preferably 5°C/s or more and more preferably 8°C/s or more.
Meanwhile, at an excessively large average cooling rate CR7, the sheet shape is degraded; thus, the average cooling rate CR7 is 100°C/s or less. The average cooling rate CR7 is preferably 50°C/s or less.
Furthermore, in order to obtain the specified retained austenite amount, the temperature T1 (cooling stop temperature T1) is 200°C or higher. The temperature T1 (cooling stop temperature T1) is preferably 220°C or higher and more preferably 240°C or higher. At a cooling stop temperature T1 exceeding (Ms temperature (°C) - 30°C), a large amount of massive non-transformed austenite remains, the amount of fresh martensite in the final cooling increases, and the stretch flangeability is degraded. Thus, the cooling stop temperature T1 is (Ms temperature (°C) - 30°C) or lower. The cooling stop temperature T1 is preferably (Ms temperature (°C) - 35°C) or lower.
Here, the average cooling rate CR7 is "(retention stop temperature T5 (cooling start temperature) (°C) - (Ms temperature (°C) - 30°C) (cooling stop temperature (°C)))/(cooling time (s) from retention stop temperature T5 to cooling stop temperature)".
Next, a member and a production method therefor according to the present invention are described.
A member of the present invention is obtained by subjecting the steel sheet of the present invention to at least one of a forming process or a joining process. Furthermore, a method for producing a member according to the present invention includes a step of subjecting the steel sheet of the present invention to at least one of a forming process or a joining process to obtain a member.
The steel sheet of the present invention has a tensile strength of 980 MPa or more and has high ductility, excellent stretch flangeability, and excellent axial crash properties. Thus, the member obtained by using the steel sheet of the present invention also has a high strength and has high ductility, excellent stretch flangeability, and excellent axial crash properties compared with typical high-strength members. Furthermore, weight reduction is possible by using the member of the present invention. Thus, the member of the present invention is suitable for use in automotive frame parts. The member of the present invention includes a welded joint.
The forming process may involve a common processing method such as pressing without any limitation. The joining process may involve a common welding such as spot welding or arc welding, riveting, crimping, or the like without any limitation.

EXAMPLES

Cold rolled steel sheets having a thickness of 1.4 mm and chemical compositions indicated in Table 1 were treated under the annealing conditions indicated in Tables 2 and 3 to produce steel sheets of the present invention and steel sheets of comparative examples.
The cold rolled steel sheets were obtained by subjecting steel slabs having chemical compositions indicated in Table 1 to hot rolling (heating temperature: 1250°C, soaking time: 60 min, finish rolling delivery temperature: 1150°C, coiling temperature: 550°C) and cold rolling (rolling reduction (rolling cumulative reduction): 50%).
Table 2 indicates conditions that do not involve the retention treatment, and Table 3 indicates conditions that involve the retention treatment.
For some of the steel sheets (cold rolled steel sheets: CR), a hot-dip galvanizing treatment was performed in one of the steps after heating from the temperature T1 to holding a temperature of 150°C or higher and 350°C or lower to prepare hot-dip galvanized steel sheets (GI).
Here, the steel sheet was dipped in a zinc coating bath at 440°C or higher and 500°C or lower to perform a hot-dip galvanizing treatment, and then the coating weight was adjusted by gas wiping or the like. The hot-dip galvanization used a zinc coating bath having an Al content of 0.10% or more and 0.22% or less. Furthermore, some of the hot-dip galvanized steel sheets were alloyed after the hot-dip galvanizing treatment into hot-dip galvannealed steel sheets (GA). Here, the alloying treatment was conducted in the temperature range of 460°C or higher and 590°C or lower. Furthermore, some steel sheets (cold rolled steel sheets: CR) were electroplated into electrogalvanized steel sheets (EG).

[Table 2]

No.	Steel No.	Heat treatment conditions											Coating conditions					Ms (°C)	Note
		Annealing temperature	Average cooling rate CR1	Temperature T1	Average heating rate	Temperature T2	Average cooling rate CR2	Temperature T3	Average cooling rate CR3	Temperature T4	Holding time	Average cooling rate CR4	Coating type	Coating bath temperature	Al amount	Alloying temperature	Timing of coating (*1)
		(°C)	(°C/s)	(°C)	(°C/s)	(°C)	(°C/s)	(°C)	(°C/s)	(°C)	(sec)	(°C/s)		(°C)	(%)	(°C)	Timing of coating (*1)
1	A	800	30	230	5	400	5	350	0.2	290	600	5	CR	-	-	-	-	332	Example
2	A	880	30	270	5	400	2	320	0.2	290	670	5	CR	-	-	-	-	376	Example
3	A	780	30	220	5	400	2	320	0.3	295	650	5	CR	-	-	-	-	304	Comparative Example
4	A	910	30	210	5	350	5	310	0.3	280	670	5	CR	-	-	-	-	381	Comparative Example
5	A	800	50	220	40	445	0.3	340	0.2	285	650	50	CR	-	-	-	-	332	Comparative Example
6	B	830	5	250	10	380	5	300	0.3	255	500	5	GA	450	0.20	500	I	364	Example
7	B	860	95	230	10	440	5	350	1	290	100	10	GA	450	0.20	500	II	403	Example
8	B	825	20	290	10	440	15	380	1	270	200	10	GA	450	0.20	500	II	356	Example
9	C	820	35	220	25	400	15	300	1	250	600	15	GA	480	0.20	500	II	312	Example
10	C	815	40	180	30	360	5	300	5	295	30	15	CR	-	-	-	-	303	Comparative Example
11	C	810	20	280	30	350	10	290	0.5	240	100	30	CR	-	-	-	-	299	Comparative Example
12	C	810	30	230	5	400	1	360	0.2	340	100	5	GA	460	0.15	550	I	299	Example
13	D	850	20	205	30	350	5	270	0.5	200	400	30	CR	-	-	-	-	250	Comparative Example
14	E	860	20	300	30	340	5	300	0.5	280	380	30	CR	-	-	-	-	419	Comparative Example
15	F	810	35	270	10	310	5	270	2	250	600	30	CR	-	-	-	-	377	Example
16	F	805	10	270	10	350	2	250	0.5	180	900	25	GA	460	0.15	480	I	368	Example
17	F	835	50	250	0.5	330	8	300	2	300	450	30	GI	460	0.15	-	III	402	Comparative Example
18	F	810	40	230	10	400	3	380	4	340	200	30	CR	-	-	-	-	374	Comparative Example
19	F	860	50	225	40	285	10	220	2	210	500	50	GI	450	0.15	-	II	411	Comparative Example
20	F	815	45	260	40	460	15	400	5	275	500	50	GI	480	0.15	-	II	381	Comparative Example
21	G	820	40	260	40	400	1	340	0.5	290	300	50	GI	480	0.15	-	II	354	Example
22	G	820	30	250	10	440	2	300	0.2	280	200	10	EG	470	0.15	-	II	354	Example
23	G	815	10	250	40	400	1	230	0.3	200	100	10	EG	470	0.15	-	III	340	Comparative Example
24	G	815	20	250	30	420	1	300	0.01	250	50	10	CR	-	-	-	-	340	Comparative Example
25	H	850	15	270	40	435	1	350	2	300	800	50	CR	-	-	-	-	333	Comparative Example
26	I	770	75	240	40	435	5	280	1	210	600	10	CR	-	-	-	-	338	Comparative Example
27	J	850	80	260	50	430	5	300	1	160	400	10	EG	470	0.13	-	II	416	Example
28	J	790	30	250	5	440	2	370	1	300	200	2	EG	480	0.13	-	II	317	Example
29	J	800	10	260	70	410	3	300	1	180	1200	2	CR	-	-	-	-	340	Comparative Example
30	J	800	1	260	50	400	3	270	2	250	200	4	CR	-	-	-	-	340	Comparative Example
31	J	795	30	260	50	430	5	300	3	130	100	10	EG	490	0.18	-	I	325	Comparative Example
32	J	805	10	240	30	440	1	400	0.2	370	200	10	EG	490	0.18	-	II	310	Comparative Example
33	K	830	40	205	50	420	3	290	2	200	300	10	CR	-	-	-	-	372	Example
34	K	840	20	240	30	420	3	320	1	300	50	10	CR	-	-	-	-	405	Example
35	K	820	30	215	50	420	3	300	1	200	10	10	CR	-	-	-	-	354	Comparative Example
36	F	805	10	240	30	330	2	280	2	200	500	0.1	GA	480	0.15	500	I	365	Comparative Example
37	L	830	10	210	70	410	3	270	0.5	170	550	20	CR	-	-	-	-	378	Comparative Example
38	M	820	40	270	70	410	3	270	0.5	155	850	20	CR	-	-	-	-	359	Comparative Example
39	N	810	50	260	50	400	10	300	5	200	600	30	CR	-	-	-	-	358	Comparative Example
40	O	895	20	300	80	400	10	300	5	200	500	30	CR	-	-	-	-	385	Comparative Example

(*1) Timing of coating: I: Immediately after heating to temperature T1, II: Immediately before holding the temperature range of 150 to 350°C, III: Immediately after holding the temperature range of 150 to 350°C.

The steel microstructure was measured by the aforementioned method. The measurement results are shown in Tables 4 and 5. Table 4 indicates conditions that do not involve the retention treatment, and Table 5 indicates conditions that involve the retention treatment.

(Tensile strength, ductility)

From each of the obtained steel sheets, a JIS No. 5 test piece for tensile test and a test piece for hole expansion test were taken and subjected to a tensile test (according to JIS Z 2241 (2011)). The tensile strength TS and the total elongation T-El are indicated in Tables 4 and 5.
Samples with a tensile strength of 980 MPa or more were determined as having excellent strength.
Furthermore, for the total elongation T-El, samples having a T-El of 14.0% or more with TS of less than 1180 MPa, samples having a T-El of 12.0% or more with TS of 1180 MPa or more and less than 1320 MPa, and samples having a T-El of 10.0% or more with TS of 1320 MPa or more were determined as having excellent ductility.

(Stretch flangeability)

The stretch flangeability was evaluated by a hole expansion test according to the Japan Iron and Steel Federation Standard, JFS T1001. That is, after a 100 mm × 100 mm square sample was punched with a punching tool having a punch diameter of 10 mm and a die diameter of 10.3 mm (13% clearance), the sample was arranged such that the burr, which had been formed by making the hole by punching, to face outward, and the hole was expanded with a conical punch having a 60 degree apex angle until cracks that penetrated through the thickness occurred. The hole expansion ratio was determined as hole expansion ratio λ (%) = {(d - d₀)/d₀} × 100 where d₀ was the initial hole diameter (mm) and d was the hole diameter (mm) at the time cracks occurred during this process. The hole expansion ratios λ are indicated in Tables 4 and 5.
Steel sheets having λ of 45% or more were determined as having excellent stretch flangeability.

(Axial crash properties)

Fig. 4 is a diagram illustrating a method for calculating the VDA bend angle.

As illustrated in Fig. 4, the VDA bend angle α was evaluated as the evaluation of the axial crash properties by a bending test (Verband der Automobilindustrie: VDA bending test) according to VDA standards (VDA238-100) standardized by the German Association of the Automotive Industry. The VDA bending test is a bending test (Verband der Automobilindustrie: VDA bending test) according to VDA standards (VDA238-100) standardized by the German Association of the Automotive Industry, and is a 3-point bending test characterized by rollers 10 with an extremely narrow roller spacing and a sharp punch 11. In the VDA bending test, a 60 mm × 60 mm square test specimen was used, the bending ridge direction was set to be parallel to the rolling direction, and the test specimen was supported by rollers with a roller diameter D of 30 mm and a roller-to-roller distance L (not illustrated) of (thickness a₀ × 2) + 0.5 mm and was pushed from above with a punch having a tip r of 0.4 mm at a stroke rate of 20 mm/min (in the drawing, reference sign 11A denotes the punch before pushing and reference sign 11B denotes the punch after pushing). Here, in Fig. 4, for the sake of convenience, only one side of the test specimen (steel sheets X_A and X_B) and the roller 10 is illustrated since this can be illustrated line-symmetrically with respect to the push axis as the center.
The VDA bend angle α is an angle (°) calculated by using equations (1) to (5) from the stroke S (mm) under the maximum load in the aforementioned bending test, and can be used as an indicator of the axial crash properties.
Steel sheets having α of 80° or more were determined as having excellent axial crash properties.

c (mm) = D / 2 + r + a_{0}

p (mm) = D / 2 + L / 2

[Math. 4]

W (mm) = \sqrt{p^{2} + {(S - c)}^{2} - c^{2}}

[Math. 5]

\sin \frac{αc}{2} = \frac{p \times c + W \times (S - c)}{p^{2} + {(S - c)}^{2}}

[Math. 6]

c o s \frac{α c}{2} = \frac{W \times p - c \times (S - c)}{p^{2} + {(S - c)}^{2}}

[Table 4]

No.	Steel No.	Microstructure										Mechanical properties				Note
		Ferrite area fraction	Tempered martensite area fraction	Lower bainite area fraction	Upper bainite area fraction	Bainite area fraction	Fresh martensite + retained γ area fraction	Fresh martensite area fraction	Remaining γ area fraction	Average C concentration in retained γ	Fe content as Fe₃C (*1)	TS	T-EI	λ	α
		(%)	(%)	(%)	(%)	(%)	(%)	(%)	(%)	(mass%)	(mass%)	(MPa)	(%)	(%)	(°)
1	A	4	61	8	11	19	16	7	9	0.81	0.07	1225	13.1	58	86	Example
2	A	0	76	10	6	16	8	2	6	0.85	0.09	1371	10.7	62	85	Example
3	A	16	43	10	7	17	24	10	14	1.04	0.06	1195	14.9	33	72	Comparative Example
4	A	0	86	7	2	9	5	1	4	0.78	0.11	1379	9.1	65	84	Comparative Example
5	A	5	64	5	10	15	16	6	10	0.80	0.24	1200	13.5	57	79	Comparative Example
6	B	6	64	6	10	16	14	5	9	0.75	0.02	996	16.7	56	87	Example
7	B	0	78	4	10	14	8	2	6	0.84	0.02	1005	14.7	59	94	Example
8	B	12	49	4	14	18	21	9	12	0.77	0.03	1120	15.7	49	81	Example
9	C	1	65	7	11	18	16	6	10	0.92	0.02	1340	12.0	50	92	Example
10	C	2	77	4	12	16	5	1	4	0.87	0.04	1379	8.7	68	88	Comparative Example
11	C	5	47	4	16	20	28	22	6	0.86	0.03	1366	9.0	30	86	Comparative Example
12	C	9	55	5	12	15	21	9	12	0.90	0.01	1322	12.4	48	81	Example
13	D	1	58	5	11	16	25	17	8	0.87	0.27	1485	9.4	39	76	Comparative Example
14	E	7	68	8	10	18	7	4	3	0.54	0.09	723	15.9	51	88	Comparative Example
15	F	6	53	11	15	26	15	9	6	0.62	0.14	1020	14.1	48	82	Example
16	F	8	51	10	20	30	11	3	8	0.61	0.07	1003	15.0	46	82	Example
17	F	0	63	16	15	31	6	1	5	0.64	0.22	1075	11.1	64	78	Comparative Example
18	F	7	53	14	19	33	7	2	5	0.74	0.34	1190	123	49	70	Comparative Example
19	F	0	70	15	10	25	5	1	4	0.46	0.11	1254	9.9	40	93	Comparative Example
20	F	3	48	15	17	32	17	11	6	0.67	0.27	1003	13.4	40	75	Comparative Example
21	G	5	67	5	8	13	15	4	11	0.89	0.03	1281	13.9	53	84	Example
22	G	5	69	4	10	14	12	3	9	0.90	0.02	1299	12.7	60	87	Example
23	G	8	63	4	13	17	12	8	4	0.59	0.01	1203	11.7	48	82	Comparative Example
24	G	7	63	5	13	18	12	6	6	0.63	0.26	1211	12.3	50	75	Comparative Example
25	H	6	31	11	15	26	37	27	10	1.04	0.01	1299	11.1	21	88	Comparative Example
26	I	1	61	9	18	27	11	9	2	0.48	0.39	1337	8.5	50	74	Comparative Example
27	J	0	58	10	20	30	12	5	7	0.62	0.14	1055	14.1	66	90	Example
28	J	8	47	5	24	29	16	8	8	0.72	0.08	987	15.4	49	82	Example
29	J	8	52	12	16	28	12	6	6	0.71	0.35	1277	12.7	52	75	Comparative Example
30	J	29	26	7	15	22	23	13	10	0.90	0.04	960	13.7	28	62	Comparative Example
31	J	10	50	7	20	27	13	9	4	0.47	0.07	1245	11.4	47	68	Comparative Example
32	J	3	74	5	10	15	8	4	4	0.68	0.31	1226	11.0	59	69	Comparative Example
33	K	1	65	6	8	14	20	5	15	0.87	0.05	1284	15.7	60	91	Example
34	K	0	72	5	6	11	17	4	13	0.88	0.01	1290	14.9	57	90	Example
35	K	8	65	2	8	10	17	9	8	0.37	0.01	1396	8.1	33	86	Comparative Example
36	F	8	49	12	21	33	10	5	5	0.70	0.24	974	13.7	46	81	Comparative Example
37	L	0	91	2	2	4	5	1	4	0.67	0.06	1346	9.4	55	82	Comparative Example
38	M	2	25	21	33	54	19	15	4	0.91	0.12	845	13.1	36	81	Comparative Example
39	N	2	64	9	10	19	15	8	7	0.72	0.15	1380	10.9	29	83	Comparative Example
40	O	7	50	8	16	24	19	5	14	0.75	0.04	1218	9.2	46	82	Comparative Example

(*1)Fe content as Fe₃C: The average content of Fe element present as carbides in tempered martensite in a region within 100 µm from the steel sheet surface in the thickness direction.

[Table 5]

No.	Steel No.	Microstructure										Mechanical properties				Note
		Ferrite area fraction	Tempered martensite area fraction	Lower bainite area fraction	Upper bainite area fraction	Bainite area fraction	Fresh martensite + retained γ area fraction	Fresh martensite area fraction	Remaining γ area fraction	Average C concentration in retained γ	Fe content as Fe₃C (*1)	TS	T-EI	λ	α
		(%)	(%)	(%)	(%)	(%)	(%)	(%)	(%)	(mass%)	(mass%)	(MPa)	(%)	(%)	(°)
41	A	5	56	10	16	26	13	6	7	0.77	0.09	1234	13.7	56	83	Example
42	A	2	61	15	12	27	10	4	6	0.69	0.05	1266	12.1	62	90	Example
43	B	1	62	5	15	20	17	7	10	0.88	0.04	1075	16.9	50	88	Example
44	G	5	62	10	10	20	13	3	10	0.84	0.03	1290	13.9	47	87	Example
45	G	3	57	5	20	25	15	4	11	0.91	0.02	1204	14.9	50	90	Example
46	G	0	69	4	14	18	13	4	9	0.90	0.01	1253	14.2	48	81	Example
47	K	0	73	3	8	11	16	5	11	0.71	0.01	1220	16.0	67	93	Example

The examples of the present invention indicated in Tables 4 and 5 are excellent in all of strength, ductility, stretch flangeability, and axial crash properties, but the comparative examples are poor in at least one.
Furthermore, it was found that a member obtained by forming a steel sheet of the example of the present invention, a member obtained by joining the steel sheet, and a member obtained by forming and joining the steel sheet had high strength, high ductility, excellent stretch flangeability, and excellent axial crash properties as with the steel sheet of the example of the present invention since the steel sheet of the example of the present invention had high strength, high ductility, excellent stretch flangeability, and excellent axial crash properties.

Industrial Applicability

The present invention enables production of a steel sheet that has excellent ductility, stretch flangeability, and axial crash properties, that is used in members of automobiles and home electric appliances, and that is particularly suitable for use in energy-absorbing members of automobiles, and such a steel sheet is suitable for press-forming of these members.

Reference Signs List

F ferrite
TM tempered martensite
UB upper bainite
LB lower bainite
FM fresh martensite
RA retained austenite
α VDA bend angle
10 roller
D roller diameter
11 punch
11A punch before pushing
11B punch after pushing
S stroke
r radius of curvature at tip
X_A steel sheet before pushing
X_B steel sheet after pushing
a₀ thickness

Claims

A steel sheet comprising:
a chemical composition containing, in mass%,
C: 0.10% or more and 0.30% or less,

Si: 0.5% or more and 2.0% or less,

Mn: 1.5% or more and 3.0% or less,

P: 0.10% or less,

S: 0.020% or less,

sol. Al: 1.00% or less,

N: 0.015% or less, and

the balance being Fe and incidental impurities; and

a steel microstructure containing, in area fractions,
ferrite: 15% or less,

tempered martensite: 30% or more and 80% or less,

bainite: 10% or more and 40% or less,

retained austenite: 5% or more, and

fresh martensite: 10% or less,

wherein:
an average C concentration in the retained austenite is 0.60 mass% or more, and

in a region within 100 µm from a steel sheet surface in a thickness direction, a content of Fe element present as carbides in the tempered martensite is 0.20 mass% or less on average.
The steel sheet according to Claim 1, wherein the chemical composition further contains, in mass%, at least one selected from
Ti: 0.100% or less,

Nb: 0.100% or less,

V: 0.100% or less,

B: 0.0050% or less,

Cr: 1.000% or less,

Cu: 1.000% or less,

Sb: 0.200% or less,

Sn: 0.200% or less,

Ta: 0.100% or less,

Ca: 0.0050% or less,

Mg: 0.0050% or less, and

REM: 0.0050% or less.
The steel sheet according to Claim 1 or 2, comprising a zinc coating layer on the steel sheet surface.
A member comprising the steel sheet according to any one of Claims 1 to 3.
A method for producing a steel sheet, the method comprising:
hot-rolling and cold-rolling a steel slab having the chemical composition according to Claim 1 or 2 to obtain a cold rolled steel sheet;

then subjecting the cold rolled steel sheet to an annealing treatment at an annealing temperature of (A_c3 temperature (°C) - 15°C) or higher and 900°C or lower;

then cooling the cold rolled steel sheet in a range from the annealing temperature to a temperature T1 of 200°C or higher and (Ms temperature (°C) - 30°C) or lower at an average cooling rate CR1 of 3°C/s or more and 100°C/s or less,

heating the cold rolled steel sheet in a temperature range from the temperature T1 to a temperature T2 of 300°C or higher and 450°C or lower at an average heating rate of 2°C/s or more;

cooling the cold rolled steel sheet in a temperature range from the temperature T2 to a temperature T3 of (T2 - 150°C) or higher and (T2 - 30°C) or lower at an average cooling rate CR2 of 0.5°C/s or more;

cooling the cold rolled steel sheet in a temperature range from the temperature T3 to a temperature T4 of 150°C or higher and 350°C or lower and the temperature T3 or lower at an average cooling rate CR3 of 0.1°C/s or more,

holding the cold rolled steel sheet in a temperature range from 150°C or higher and 350°C or lower and the temperature T4 or lower for 20 s or longer and 1000 s or shorter; and

cooling the cold rolled steel sheet to a temperature of 50°C or lower at an average cooling rate CR4 of 1°C/s or more.
A method for producing a steel sheet, the method comprising:
hot-rolling and cold-rolling a steel slab having the chemical composition according to Claim 1 or 2 to obtain a cold rolled steel sheet;

then subjecting the cold rolled steel sheet to an annealing treatment at an annealing temperature of (A_c3 temperature (°C) - 15°C) or higher and 900°C or lower;

then cooling the cold rolled steel sheet in a temperature range from the annealing temperature to 500°C at an average cooling rate CR5 of 5°C/s or more and 100°C/s or less;

retaining the cold rolled steel sheet in a temperature range from 500°C to a retention stop temperature T5 of an Ms temperature (°C) or higher and 320°C or higher at an average cooling rate CR6 of 10°C/s or less for 10 s or longer and 60 s or shorter;

cooling the cold rolled steel sheet in a temperature range from the retention stop temperature T5 to a temperature T1 of 200°C or higher and (Ms temperature (°C) - 30°C) or lower at an average cooling rate CR7 of 3°C/s or more and 100°C/s or less;

heating the cold rolled steel sheet in a temperature range from the temperature T1 to a temperature T2 of 300°C or higher and 450°C or lower at an average heating rate of 2°C/s or more;

cooling the cold rolled steel sheet in a temperature range from the temperature T2 to a temperature T3 of (T2 - 150°C) or higher and (T2 - 30°C) or lower at an average cooling rate CR2 of 0.5°C/s or more;

cooling the cold rolled steel sheet in a temperature range from the temperature T3 to a temperature T4 of 150°C or higher and 350°C or lower and the temperature T3 or lower at an average cooling rate CR3 of 0.1°C/s or more,

holding the cold rolled steel sheet in a temperature range from 150°C or higher and 350°C or lower and the temperature T4 or lower for 20 s or longer and 1000 s or shorter; and

cooling the cold rolled steel sheet to a temperature of 50°C or lower at an average cooling rate CR4 of 1°C/s or more.
The method for producing a steel sheet according to Claim 5, further comprising hot-dip-galvanizing or hot-dip-galvannealing the steel sheet after the steel sheet is cooled at the average cooling rate CR1 of 3°C/s or more and 100°C/s or less.
The method for producing a steel sheet according to Claim 6, further comprising hot-dip-galvanizing or hot-dip-galvannealing the steel sheet after the steel sheet is cooled at the average cooling rate CR7 of 3°C/s or more and 100°C/s or less.
The method for producing a steel sheet according to Claim 5 or 6, further comprising electrogalvanizing the steel sheet after the steel sheet is cooled at the average cooling rate CR4 of 1°C/s or more.
A method for producing a member, the method comprising a step of preparing a member by subjecting the steel sheet according to any one of Claims 1 to 3 to either or both of a forming process and a joining process.