Steel plates and parts containing them

Abstract

A steel sheet is provided that has a specific chemical composition, comprising, by area%, 20-70% soft ferrite, 5-35% retained austenite, 8-35% fresh martensite, and 0-60% total tempered martensite and bainite, a microstructure in which the integral intensity / peak height ratio of the X-ray diffraction peak on the (220) plane of the fcc phase at a position 1 / 4 of the sheet thickness from the surface is 0.100 deg. or less, and the integral intensity / peak height ratio of the X-ray diffraction peak on the (311) plane of the fcc phase at a position 1 / 4 of the sheet thickness from the surface is 0.120 deg. or less, and a Vickers hardness of 300 Hv or more.

Description

Technical Field

[0001] The present invention relates to steel plates and components containing the same.

Background Art

[0002] In recent years, from the perspective of regulations on greenhouse gas emissions associated with measures against global warming, improvement of the fuel efficiency of automobiles has been demanded. For the sake of achieving both vehicle body weight reduction and collision safety, the application of high-strength steel plates in automobile parts has been expanding. However, for steel plates used in automobile parts, not only strength but also various workabilities such as press formability and weldability are required in part forming.

[0003] In relation to this, for example, in Patent Document 1, a high-strength and high-ductility steel plate having a predetermined component composition, in which the steel structure has, in terms of the area ratio to the entire structure, retained austenite of 8% or more and ferrite of more than 5% and 50% or less, and the balance is composed of one or more of bainite, martensite, tempered bainite, and tempered martensite, and regarding the carbon concentration in the retained austenite, the average carbon concentration is 0.8 to 1.1% by mass, the standard deviation of the carbon concentration distribution is 0.25% by mass or more, and the area ratio of the region where the carbon concentration is 1.3% by mass or more to the entire structure is 1.0% or more is described. Further, in Patent Document 1, according to the above configuration, not only the amount (area ratio) of retained austenite and the average carbon concentration are defined, but also the distribution of the carbon concentration is controlled, so that the TRIP phenomenon is developed from the initial stage to the later stage of deformation, and a high work hardening rate is realized, whereby a high-strength and high-ductility steel plate excellent in strength-ductility balance, having a tensile strength (TS) of 980 MPa or more and ensuring a tensile strength (TS) × elongation (EL) of 25000 MPa·% or more, can be provided.

[0004] Patent Document 2 defines a steel structure in the range of 1 / 8 to 3 / 8 thickness centered at a position 1 / 4 thickness from the surface as containing, by volume fraction, soft ferrite: 0% to 30%, retained austenite: 3% to 40%, fresh martensite: 0% to 30%, and the sum of pearlite and cementite: 0% to 10%, with the remainder being hard ferrite, and in the range of 1 / 8 to 3 / 8 thickness centered at the position 1 / 4 thickness from the surface, the proportion of retained austenite with an aspect ratio of 2.0 or more to all retained austenite is 50% or more, and the region having a hardness of 80% or less of the hardness of the range of 1 / 8 to 3 / 8 thickness is defined as the soft layer. The invention describes a steel sheet characterized in that, when prepared, a soft layer with a thickness of 1 to 100 μm exists in the thickness direction from the surface, the volume fraction of ferrite crystal grains with an aspect ratio of less than 3.0 is 50% or more of the ferrite crystal grains contained in the soft layer, the volume fraction of retained austenite in the soft layer is 50% or more of the volume fraction of retained austenite in the range of 1 / 8 thickness to 3 / 8 thickness, and when the emission intensity of the wavelength indicating Si is analyzed by high-frequency glow discharge analysis from the surface in the thickness direction, a peak of the emission intensity of the wavelength indicating Si appears in the range of more than 0.2 μm from the surface and 5 μm or less from the surface. Furthermore, Patent Document 2 states that according to the above configuration, a steel sheet can be provided that has good ductility and hole-expanding properties, as well as excellent fatigue resistance, bendability, and excellent plating adhesion. [Prior art documents] [Patent Documents]

[0005] [Patent Document 1] Japanese Patent Publication No. 2016-098413 [Patent Document 2] International Publication No. 2019 / 186997 [Overview of the Initiative] [Problems that the invention aims to solve]

[0006] As steel sheets become stronger, their workability generally decreases, and properties such as ductility, as described in Patent Documents 1 and 2, are known to decline. Furthermore, automotive steel sheets are often processed into the desired part shape by press forming. Typically, press forming is carried out in multiple steps, resulting in a relatively large number of areas where, for example, strain accumulates within the steel sheet after primary deformation, and then undergoes further deformation. However, since steel sheets work harden and become stronger when strain is introduced, their workability in subsequent processes generally decreases. Therefore, steel sheets are required to exhibit high formability even when a certain degree of strain has been introduced, for example, by possessing excellent work hardening ability (the ability to continue to harden).

[0007] Therefore, the present invention aims to provide a steel sheet that, despite having high strength, exhibits improved ductility and work hardening ability in the later stages of deformation, through a novel configuration. [Means for solving the problem]

[0008] To achieve the above objectives, the inventors focused on the chemical composition of steel sheets and specific microstructures. Specifically, the inventors first found that by appropriately controlling the chemical composition of the steel sheet, and further controlling the microstructure so that the Vickers hardness is 300 Hv or higher, and by constructing the microstructure with a predetermined amount of retained austenite, it is possible to improve the ductility of the steel sheet while achieving high strength. In addition, the inventors found that by optimizing the relationship between the Si, Mn, and Al content in the steel sheet and appropriately controlling the morphology of retained austenite observed by X-ray diffraction, it is possible to further improve the ductility of the steel sheet and significantly improve the work hardening ability in the later stages of deformation, thus completing the present invention.

[0009] The present invention, which has achieved the above objectives, is as follows. (1) The chemical composition is expressed in mass%, C: 0.050%~0.500%, Si: 0.20%~2.00% Mn: 1.00%~4.00% P: 0.1000% or less, S: 0.0100% or less, Al: 0.200%~1.500% N: 0.0150% or less, O: 0.0100% or less, Ti: 0~0.300%, V: 0~1.00%, Nb: 0~0.100%, Cr: 0~2.00%, Ni: 0~2,000%, Cu: 0~2.000%, Co: 0~2.00%, Mo: 0~1.00%, W: 0~1.00%, B: 0~0.0100%, Ta: 0~1.00%, Sn: 0~1.00%, Sb: 0~1.00%, Ca: 0~0.0500%, Mg: 0~0.0500%, Zr: 0~0.5000%, Hf: 0~0.0100%, Bi: 0~0.0100%, REM: 0~0.0100%, As: 0~0.100%, and The remainder consists of Fe and impurities, and satisfies the following formula 1. 1.00≦Mn / (Al+0.5Si)≦4.00...Equation 1 The metallic structure, in area percentage, Soft ferrite: 20-70%, Residual austenite: 5-35%, Fresh martensite: 8-35%, and Total of tempered martensite and bainite: 0-60%, The ratio of the integrated intensity / peak height of the X-ray diffraction peak in the (220) plane of the fcc phase at a position 1 / 4 of the plate thickness from the surface is 0.100 deg. or less, and The ratio of the integrated intensity to the peak height of the X-ray diffraction peak on the (311) plane of the fcc phase at a position of 1 / 4 of the plate thickness from the surface is 0.120 deg. or less, A steel sheet characterized by having a Vickers hardness of 300 Hv or more. (2) The chemical composition is in mass%, N: 0.0100% or less, Ca: 0 to 0.0100%, Mg: 0 to 0.0100%, and Zr: 0 to 0.0100% The steel sheet according to (1) above, characterized by containing the same. (3) The chemical composition is in mass%, Ti: 0.001 to 0.300%, V: 0.001 to 1.00%, Nb: 0.001 to 0.100%, Cr: 0.001 to 2.00%, Ni: 0.001 to 2.000%, Cu: 0.001 to 2.000%, Co: 0.001 to 2.00%, Mo: 0.001 to 1.00%, W: 0.001 to 1.00%, B: 0.0001 to 0.0100%, Ta: 0.0,01 to 1.00%, Sn: 0.001 to 1.00%, Sb: 0.001 to 1.00%, Ca: 0.0001 to 0.0500%, Mg: 0.0001 to 0.0500%, Zr: 0.0001 to 0.5000%, Hf: 0.0001 to 0.0100%, Bi: 0.0001 to 0.0100%, and REM: 0.0001 to 0.0100%, As: 0.001 to 0.100% The steel sheet according to (1) or (2) above, characterized by containing at least one of the same. (4) The metal structure is in area%, A steel sheet according to any one of the above items (1) to (3), characterized by containing 12 to 35% fresh martensite. (5) A component characterized by including a steel plate as described in any one of the above items (1) to (4). [Effects of the Invention]

[0010] According to the present invention, it is possible to provide a steel sheet that has high strength while also having improved ductility and work hardening ability in the later stages of deformation. [Modes for carrying out the invention]

[0011] <Steel plate> The steel sheet according to the embodiment of the present invention has a chemical composition in mass%, C: 0.050%~0.500%, Si: 0.20%~2.00% Mn: 1.00%~4.00% P: 0.1000% or less, S: 0.0100% or less, Al: 0.200%~1.500% N: 0.0150% or less, O: 0.0100% or less, Ti: 0~0.300%, V: 0~1.00%, Nb: 0~0.100%, Cr: 0~2.00%, Ni: 0~2,000%, Cu: 0~2.000%, Co: 0~2.00%, Mo: 0~1.00%, W: 0~1.00%, B: 0~0.0100%, Ta: 0~1.00%, Sn: 0~1.00%, Sb: 0~1.00%, Ca: 0~0.0500%, Mg: 0~0.0500%, Zr: 0~0.5000%, Hf: 0~0.0100%, Bi: 0~0.0100%, REM: 0~0.0100%, As: 0~0.100%, and The remainder consists of Fe and impurities, and satisfies the following formula 1. 1.00≦Mn / (Al+0.5Si)≦4.00...Equation 1 The metallic structure, in area percentage, Soft ferrite: 20-70%, Residual austenite: 5-35%, Fresh martensite: 8-35%, and Total of tempered martensite and bainite: 0-60%, The ratio of the integrated intensity / peak height of the X-ray diffraction peak in the (220) plane of the fcc phase at a position 1 / 4 of the plate thickness from the surface is 0.100 deg. or less, and The ratio of the integrated intensity / peak height of the X-ray diffraction peak in the (311) plane of the fcc phase at a position 1 / 4 of the plate thickness from the surface is 0.120 deg. or less. It is characterized by having a Vickers hardness of 300 Hv or higher.

[0012] As mentioned earlier, it is generally known that increasing the strength of a steel sheet reduces its workability and other properties such as ductility. Therefore, the present inventors first found that by appropriately controlling the chemical composition of the steel sheet, particularly by limiting the Si content to 2.00 mass% or less while increasing the Al content to 0.200 mass% or more, and further controlling the microstructure so that the Vickers hardness is 300 Hv or more, and by configuring the microstructure to contain a predetermined amount of retained austenite, more specifically by configuring the microstructure of the steel sheet to contain retained austenite: 5-35% by area%, it is possible to improve the ductility of the steel sheet while achieving high strength.

[0013] More specifically, while Si can improve the ductility of steel sheets by generating retained austenite and exhibiting the TRIP (transformation-induced plasticity) effect, excessive Si content in the steel can reduce LME resistance. On the other hand, Al has the effect of increasing retained austenite by suppressing the formation of iron-based carbides. Therefore, the inventors have found that by including a relatively large amount of Al in the steel, which has a similar effect in generating retained austenite, in place of some of the Si, and more specifically by limiting the Si content to 2.00 mass% or less while including 0.200 mass% or more of Al, the LME resistance of the steel sheet can be improved, and the metal structure of the steel sheet can be configured to contain retained austenite: 5-35% by area%, thereby improving the ductility of the steel sheet, particularly its uniform elongation. Furthermore, the inventors have found that by controlling the content of other alloying elements within a predetermined range, and by appropriately controlling the microstructure other than retained austenite so that the Vickers hardness is 300 Hv or higher, it is possible to achieve not only the above-mentioned improvement in ductility, but also increased strength of the steel sheet, for example, a tensile strength of 980 MPa or higher.

[0014] Generally, it is known that including retained austenite in the metal structure increases the work hardening rate due to the TRIP effect, thereby improving properties such as uniform elongation. On the other hand, in forming operations that are divided into multiple processes, such as press forming, workability generally decreases in the later stages of deformation due to work hardening caused by strain introduced in the early stages of deformation. Therefore, there is a need for steel sheets that can achieve both high strength and workability by exhibiting high work hardening ability even in the later stages of deformation such as press forming. However, simply controlling retained austenite within a predetermined range can improve work hardening ability in the early stages of deformation, but it may not be possible to sufficiently improve work hardening ability in the later stages of deformation. Therefore, the inventors of this invention further investigated by focusing on the morphology of retained austenite contained in steel sheets. As a result, the inventors have found that by optimizing the relationship between the Si, Mn, and Al content in the steel sheet, more specifically by controlling it to satisfy 1.00 ≤ Mn / (Al + 0.5Si) ≤ 4.00, and by appropriately controlling the morphology of retained austenite observed by X-ray diffraction, more specifically by controlling the ratio of the integrated intensity (cps·deg.) / peak height (cps) of the X-ray diffraction peak at the (220) plane of the fcc phase at a position 1 / 4 of the thickness from the steel sheet surface to 0.100 (deg.) or less, and the ratio of the integrated intensity (cps·deg.) / peak height (cps) of the X-ray diffraction peak at the (311) plane of the fcc phase at a position 1 / 4 of the thickness from the steel sheet surface to 0.120 (deg.) or less, the ductility of the steel sheet can be further improved, and the work hardening ability in the later stages of deformation can be significantly improved.

[0015] While not intended to be bound by any particular theory, it is believed that by ensuring that the Si, Mn, and Al content in the steel sheet satisfies 1.00 ≤ Mn / (Al + 0.5Si) ≤ 4.00, the Ac3 point can be controlled to a relatively low value, as will be explained in more detail later in relation to the manufacturing method, and the phase transformation to austenite can be optimized. As a result, it is believed that the retained austenite in the final resulting metal structure can be controlled to be fine and needle-shaped. It is believed that by including a relatively large amount of such fine and needle-shaped retained austenite, the ductility of the steel sheet can be further improved, and the work hardening ability in the later stages of deformation can be enhanced.

[0016] More specifically, in fine, needle-shaped retained austenite (hereinafter also referred to as "fine needle-shaped retained austenite"), it is believed that alloying elements, specifically C and Mn, are present in relatively high concentrations due to their fine and needle-like shape. In particular, in fine needle-shaped retained austenite, the migration speed of the interface is small, allowing for the distribution of alloying elements such as C and Mn at the migration interface, thus enabling high concentrations of these alloying elements. Furthermore, it is believed that the presence of a larger amount of fine needle-shaped retained austenite compared to other forms of retained austenite, such as massive retained austenite, results in relatively sharp X-ray diffraction peaks, namely the (220) and (311) planes of the fcc phase. Here, since alloying elements such as C and Mn have a stabilizing effect on austenite, it is believed that fine needle-shaped retained austenite, in which these alloying elements are present in relatively high concentrations, exhibits significantly higher stability compared to ordinary retained austenite. In this case, even if the steel sheet is processed in the early stages of deformation such as press forming, it is thought that at least a portion of the retained austenite, preferably a relatively large portion of the retained austenite, will remain in the metal structure even in the later stages of deformation without transforming into a hard phase such as martensite. Based on this finding, further investigations have shown that by controlling the X-ray diffraction peaks corresponding to the needle-shaped retained austenite to have a sharper shape (i.e., controlling the ratio of the integrated intensity to the peak height of the corresponding X-ray diffraction peak to be smaller), specifically by controlling the ratio of the integrated intensity to the peak height of the X-ray diffraction peak on the (220) plane of the fcc phase at a position 1 / 4 of the thickness from the steel sheet surface to 0.100 deg. or less, and by controlling the ratio of the integrated intensity to the peak height of the X-ray diffraction peak on the (311) plane of the fcc phase at a position 1 / 4 of the thickness from the steel sheet surface to 0.120 deg. or less, it is possible to sufficiently form fine needle-shaped retained austenite with a relatively high concentration of alloying elements.As a result, the inventors have found that they can further improve the ductility of the steel sheet and significantly increase the work hardening rate in the high-strain region, thereby achieving a high work hardening rate even in conditions where some strain has been introduced, such as in the later stages of deformation during press forming. For example, in electron backscattering diffraction (EBSD), areas containing noise may be identified as retained austenite. Furthermore, because the above-mentioned fine needle-shaped retained austenite is so fine, it is difficult to obtain information about it using EBSD while reliably eliminating the influence of noise. Therefore, in the steel sheet according to the embodiment of the present invention, the morphology of retained austenite is defined using information obtained by X-ray diffraction, rather than EBSD.

[0017] Therefore, according to the steel sheet according to the embodiment of the present invention, despite having high strength due to the optimization of the chemical composition and metal structure of the steel sheet, the ductility and work hardening ability in the later stages of deformation can be sufficiently improved. For this reason, the steel sheet according to the embodiment of the present invention is particularly useful for use in the automotive sector, where a high level of both high strength and formability is required.

[0018] The steel sheets according to embodiments of the present invention will be described in more detail below. In the following description, "%", which is the unit for the content of each element, means "mass%" unless otherwise specified. In this specification, "~", which indicates a numerical range, is used to mean that the numbers written before and after it are included as the lower and upper limits, respectively, unless otherwise specified.

[0019] [C:0.050~0.500%] Carbon (C) is an essential element for ensuring the strength and hardness of steel sheets. Furthermore, because C stabilizes austenite, it is effective in stabilizing retained austenite and improving work hardening ability in the high-strain range. To fully obtain these effects, the C content should be 0.050% or higher. The C content may also be 0.060% or higher, 0.080% or higher, 0.100% or higher, 0.120% or higher, 0.140% or higher, 0.160% or higher, 0.180% or higher, 0.200% or higher, 0.220% or higher, 0.240% or higher, or 0.260% or higher. On the other hand, excessive C content may lead to a decrease in ductility due to an excessive increase in strength. Therefore, the C content should be 0.500% or lower. The C content may be 0.480% or less, 0.460% or less, 0.440% or less, 0.420% or less, 0.400% or less, 0.380% or less, 0.360% or less, 0.340% or less, 0.320% or less, 0.300% or less, or 0.280% or less.

[0020] [Si: 0.20~2.00%] Si is an element that contributes to improved strength and formability by suppressing the formation of iron-based carbides and promoting the formation of retained austenite. To fully obtain these effects, the Si content should be 0.20% or more. The Si content may be 0.30% or more, 0.40% or more, 0.50% or more, 0.60% or more, 0.80% or more, 1.00% or more, or 1.10% or more. On the other hand, excessive Si content may reduce LME resistance. Therefore, the Si content should be 2.00% or less. The Si content may be 1.90% or less, 1.80% or less, 1.70% or less, 1.60% or less, 1.50% or less, 1.40% or less, 1.30% or less, or 1.20% or less.

[0021] [Mn: 1.00~4.00%] Mn is an element that enhances hardenability and contributes to improved strength and hardness. To fully obtain these effects, the Mn content should be 1.00% or more. The Mn content may be 1.20% or more, 1.40% or more, 1.60% or more, 1.80% or more, 2.00% or more, or 2.20% or more. On the other hand, excessive Mn content may lead to excessive formation of the hard phase, which can reduce ductility. Therefore, the Mn content should be 4.00% or less. The Mn content may be 3.80% or less, 3.60% or less, 3.40% or less, 3.20% or less, 3.00% or less, 2.80% or less, 2.60% or less, or 2.40% or less.

[0022] [P:0.1000% or less] P is a solid solution strengthening element and is effective in increasing the strength of steel plates, but excessive addition may degrade weldability and toughness. Therefore, the P content should be 0.1000% or less. Preferably, the P content is 0.0500% or less, 0.0400% or less, or 0.0200% or less. The P content may be 0%, but reducing the P content to an extreme degree will increase the cost of removing P. For this reason, from an economic standpoint, the P content may be 0.0001% or more, 0.0005% or more, or 0.0010% or more.

[0023] [S:0.0100% or less] S is an element present as an impurity, and it can form MnS in steel, reducing toughness and hole-expanding properties. Therefore, the S content should be 0.0100% or less. Preferably, the S content is 0.0050% or less, 0.0040% or less, 0.0030% or less, or 0.0020% or less. The S content may be 0%, but extremely low S content increases desulfurization costs. For this reason, from an economic standpoint, the S content may be 0.0001% or more, 0.0003% or more, or 0.0005% or more.

[0024] [Al: 0.200~1.500%] Al is an element that acts as a deoxidizing agent. Furthermore, Al has the effect of increasing retained austenite by suppressing the formation of iron-based carbides. To fully obtain these effects, the Al content should be 0.200% or higher. The Al content may also be 0.250% or higher, 0.300% or higher, 0.350% or higher, 0.400% or higher, 0.450% or higher, 0.500% or higher, 0.550% or higher, or 0.600% or higher. On the other hand, excessive Al content saturates the effect, and including more Al than necessary in the steel sheet leads to increased manufacturing costs. Also, in relation to the Mn content, excessive Al content can raise the Ac3 point, potentially resulting in a coarser microstructure. Therefore, the Al content should be 1.500% or lower. The Al content may also be 1.400% or lower, 1.200% or lower, 1.000% or lower, 0.800% or lower, or 0.700% or lower.

[0025] [N:0.0150% or less] N is an element present as an impurity, and a high N content can lead to the formation of coarse nitrides in the steel, reducing its bendability and hole-expanding properties. Therefore, the N content should be 0.0150% or less. Preferably, the N content is 0.0120% or less, 0.0100% or less, 0.0080% or less, 0.0060% or less, or 0.0050% or less. The N content may be 0%, but extremely low N content increases the cost of nitrogen removal. For this reason, from an economic standpoint, the N content may be 0.0001% or more, 0.0003% or more, or 0.0005% or more.

[0026] [O:0.0100% or less] O is an element present as an impurity, and a high O content can lead to the formation of coarse oxides in the steel, reducing its bendability and hole-expanding properties. Therefore, the O content should be 0.0100% or less. Preferably, the O content is 0.0080% or less, 0.0060% or less, or 0.0050% or less. The O content may be 0%, but extremely low O content increases manufacturing costs. For this reason, from the viewpoint of manufacturing costs, the O content may be 0.0001% or more, 0.0003% or more, or 0.0005% or more.

[0027] The basic chemical composition of the steel sheet according to the embodiment of the present invention is as described above. Furthermore, the steel sheet may optionally contain at least one of the following elements in place of a portion of the remaining Fe.

[0028] [Ti: 0~0.300%] Ti is an element that contributes to increasing the strength of steel sheets through precipitation strengthening, fine grain strengthening by suppressing the growth of ferrite crystal grains, and dislocation strengthening through the suppression of recrystallization. The Ti content may be 0%, but to obtain these effects, it is preferable that the Ti content be 0.001% or more. The Ti content may be 0.005% or more, 0.010% or more, or 0.050% or more. On the other hand, if the Ti content is excessive, the precipitation of carbonitrides may increase. For this reason, it is preferable that the Ti content be 0.300% or less. The Ti content may be 0.250% or less, 0.200% or less, 0.150% or less, 0.100% or less, or 0.080% or less.

[0029] [V: 0~1.00%] V is an element that contributes to increasing the strength of steel sheets through precipitation strengthening, fine grain strengthening by suppressing the growth of ferrite crystal grains, and dislocation strengthening through the suppression of recrystallization. The V content may be 0%, but to obtain these effects, it is preferable that the V content be 0.001% or more. The V content may be 0.005% or more, 0.01% or more, or 0.02% or more. On the other hand, if the V content is excessive, the precipitation of carbonitrides may increase. For this reason, it is preferable that the V content be 1.00% or less. The V content may be 0.50% or less, 0.40% or less, 0.30% or less, or 0.10% or less.

[0030] [Nb: 0~0.100%] Nb is an effective element for controlling the morphology of carbides and is also effective in improving toughness by refining the structure. The Nb content may be 0%, but to obtain these effects, it is preferable that the Nb content be 0.001% or more. The Nb content may be 0.005% or more, 0.010% or more, or 0.020% or more. On the other hand, if the Nb content is excessive, coarse Nb carbides may be formed. For this reason, it is preferable that the Nb content be 0.100% or less. The Nb content may be 0.080% or less, 0.060% or less, 0.050% or less, or 0.040% or less.

[0031] [Cr: 0~2.00%] Cr is an element that is effective in increasing the strength of steel by improving hardenability. The Cr content may be 0%, but to obtain such an effect, it is preferable that the Cr content be 0.001% or more. The Cr content may be 0.005% or more, 0.01% or more, 0.02% or more, or 0.05% or more. On the other hand, if the Cr content is excessive, the workability at hot temperatures may be impaired, which may reduce productivity. For this reason, it is preferable that the Cr content be 2.00% or less. The Cr content may be 1.80% or less, 1.50% or less, 1.00% or less, or 0.50% or less.

[0032] [Ni: 0~2.000%] Ni is an element effective in increasing the strength of steel plates. Ni is also effective in improving wettability and promoting alloying reactions. While the Ni content may be 0%, it is preferable that the Ni content be 0.001% or higher to obtain these effects. The Ni content may be 0.010% or higher, 0.020% or higher, 0.050% or higher, 0.100% or higher, or 0.200% or higher. On the other hand, excessive Ni content may reduce weldability. Therefore, it is preferable that the Ni content be 2.000% or lower. The Ni content may be 1.500% or lower, 1.200% or lower, 1.000% or lower, 0.800% or lower, 0.600% or lower, or 0.300% or lower.

[0033] [Cu: 0~2.000%] Cu is an effective element for improving the strength of steel plates. While the Cu content may be 0%, it is preferable that the Cu content be 0.001% or more to obtain this effect. The Cu content may also be 0.010% or more, 0.100% or more, or 0.300% or more. On the other hand, excessive Cu content may reduce weldability. For this reason, it is preferable that the Cu content be 2.000% or less. The Cu content may also be 1.500% or less, 1.200% or less, 1.000% or less, 0.600% or less, or 0.400% or less.

[0034] [Co: 0~2.00%] Co is an element that is effective in increasing the hardenability and thus the strength of steel. The Co content may be 0%, but to obtain this effect, it is preferable that the Co content be 0.001% or more. The Co content may be 0.01% or more, 0.10% or more, 0.20% or more, or 0.30% or more. On the other hand, if the Co content is excessive, coarse Co carbides may precipitate. For this reason, it is preferable that the Co content be 2.00% or less. The Co content may be 1.50% or less, 1.00% or less, 0.80% or less, 0.60% or less, or 0.40% or less.

[0035] [Mo: 0~1.00%] Mo is an element that suppresses phase transformation at high temperatures and is effective in increasing strength. The Mo content may be 0%, but to obtain such an effect, it is preferable that the Mo content be 0.001% or more. The Mo content may be 0.01% or more, 0.05% or more, 0.10% or more, or 0.15% or more. On the other hand, if the Mo content is excessive, the effect may saturate or coarse intermetallic compounds and carbides may be formed. For this reason, it is preferable that the Mo content be 1.00% or less. The Mo content may be 0.80% or less, 0.60% or less, or 0.30% or less.

[0036] [W: 0~1.00%] W is an element that suppresses phase transformation at high temperatures and is effective in increasing strength. The W content may be 0%, but to obtain such an effect, it is preferable that the W content be 0.001% or more. The W content may be 0.01% or more, 0.05% or more, 0.10% or more, or 0.15% or more. On the other hand, if the W content is excessive, the workability at hot temperatures may be impaired and productivity may decrease. For this reason, it is preferable that the W content be 1.00% or less. The W content may be 0.80% or less, 0.60% or less, 0.50% or less, or 0.30% or less.

[0037] [B: 0~0.0100%] B is an element that suppresses the formation of ferrite and pearlite during the cooling process from the austenite temperature range and promotes the formation of low-temperature transformation structures such as martensite. Furthermore, B is a beneficial element for increasing the strength of steel. While the B content may be 0%, it is preferable that the B content be 0.0001% or higher to obtain these effects. The B content may also be 0.0003% or higher, 0.0005% or higher, or 0.0010% or higher. On the other hand, excessive B content may lead to the formation of coarse B oxides in the steel. Therefore, it is preferable that the B content be 0.0100% or lower. The B content may also be 0.0080% or lower, 0.0060% or lower, 0.0050% or lower, or 0.0020% or lower.

[0038] [Ta: 0~1.00%] Ta is an effective element for controlling the morphology of carbides and increasing the strength of steel sheets. While the Ta content may be 0%, it is preferable that the Ta content be 0.001% or higher to obtain these effects. The Ta content may also be 0.002% or higher, 0.005% or higher, or 0.008% or higher. On the other hand, if the Ta content is excessive, the effect will saturate, and including more Ta than necessary in the steel sheet will lead to an increase in manufacturing costs. For this reason, it is preferable that the Ta content be 1.00% or lower. The Ta content may also be 0.80% or lower, 0.60% or lower, 0.40% or lower, 0.20% or lower, or 0.10% or lower.

[0039] [Sn: 0~1.00%] [Sb: 0~1.00%] Sn and Sb are elements that are effective in improving corrosion resistance. The Sn and Sb content may be 0%, but to obtain such an effect, it is preferable that the content of each element be 0.001% or more, and may be 0.01% or more, 0.02% or more, or 0.05% or more. On the other hand, excessive content of these elements may lead to a decrease in toughness. Therefore, it is preferable that the Sn and Sb content be 1.00% or less, and may be 0.80% or less, 0.50% or less, 0.30% or less, 0.20% or less, 0.10% or less, or 0.08% or less.

[0040] [Ca: 0~0.0500%] [Mg: 0~0.0500%] [Zr:0~0.5000%] [Hf: 0~0.0100%] [REM:0~0.0100%] Ca, Mg, Zr, Hf, and REM are elements that contribute to improving the formability of steel sheets. The content of Ca, Mg, Zr, Hf, and REM may be 0%, but to obtain such an effect, it is preferable that the content of Ca, Mg, Zr, Hf, and REM be 0.0001% or more, and may be 0.0005% or more, 0.0010% or more, or 0.0015% or more. On the other hand, if these elements are present in excess, the ductility of the steel sheet may decrease. Therefore, it is preferable that the content of Ca and Mg be 0.0500% or less, and may be 0.0200% or less, 0.0100% or less, 0.0080% or less, 0.0060% or less, 0.0040% or less, or 0.0020% or less. The Zr content is preferably 0.5000% or less, and may be 0.1000% or less, 0.0100% or less, 0.0080% or less, 0.0060% or less, 0.0040% or less, or 0.0020% or less. The Hf and REM content are preferably 0.0100% or less, and may be 0.0080% or less, 0.0060% or less, 0.0040% or less, or 0.0020% or less, respectively. In this specification, REM refers to the collective term for 17 elements including scandium (Sc) with atomic number 21, yttrium (Y) with atomic number 39, and the lanthanides from lanthanum (La) with atomic number 57 to lutetium (Lu) with atomic number 71, and the REM content is the total content of these elements.

[0041] [Bi: 0~0.0100%] Bi is an effective element for improving corrosion resistance. While the Bi content may be 0%, it is preferable that the Bi content be 0.0001% or more to obtain such an effect. The Bi content may also be 0.0005% or more, 0.0010% or more, or 0.0015% or more. On the other hand, if the Bi content is excessive, the effect will saturate, and including more Bi than necessary in the steel sheet will lead to an increase in manufacturing costs. Therefore, it is preferable that the Bi content be 0.0100% or less. The Bi content may also be 0.0050% or less, or 0.0030% or less.

[0042] [As: 0~0.100%] As is an effective element for improving corrosion resistance. While the As content may be 0%, it is preferable that the As content be 0.001% or more, and may be 0.005% or more, or 0.008% or more, in order to obtain such an effect. On the other hand, if the As content is excessive, the effect will saturate, and including more As than necessary in the steel sheet will lead to an increase in manufacturing costs. Therefore, the As content should be 0.100% or less, and may be 0.050% or less, 0.030% or less, 0.020% or less, or 0.010% or less.

[0043] In the steel sheet according to the embodiment of the present invention, the remainder of the elements other than those mentioned above consists of Fe and impurities. Impurities include components that are mixed in during the industrial production of steel sheets due to various factors in the manufacturing process, such as raw materials like ore and scrap, and components that are included in a range that does not affect the effects of the present invention.

[0044] The chemical composition of the steel sheet according to the embodiment of the present invention may be measured by a general analytical method. For example, the chemical composition of the steel sheet may be measured using inductively coupled plasma-atomic emission spectrometry (ICP-AES). C and S may be measured using the combustion-infrared absorption method, N may be measured using the inert gas fusion-thermal conductivity method, and O may be measured using the inert gas fusion-nondispersive infrared absorption method. If molten steel analysis values, slab analysis values, or steel sheet analysis values ​​of other steel sheets manufactured from the same molten steel can be confirmed, the analysis of a test piece taken from the steel sheet may be omitted, and those analysis values ​​may be considered as the chemical composition of the steel sheet.

[0045] [1.00 ≤ Mn / (Al + 0.5Si) ≤ 4.00] The chemical composition of the steel sheet according to the embodiment of the present invention must satisfy the following formula 1. 1.00≦Mn / (Al+0.5Si)≦4.00...Equation 1 In the formula, Mn, Al, and Si represent the mass percentage content of each element. As mentioned earlier, by ensuring that the Si, Mn, and Al content in the steel sheet satisfies 1.00 ≤ Mn / (Al + 0.5Si) ≤ 4.00, the Ac3 point can be controlled to a relatively low value, and the phase transformation to austenite can be optimized. As a result, it becomes possible to control the retained austenite in the final resulting metal structure to be fine and needle-shaped.

[0046] As will be explained in more detail later in relation to the manufacturing method, the inventors have found that by constructing the microstructure of a steel sheet before secondary heat treatment with a structure mainly composed of martensite, and then performing secondary heat treatment on a steel sheet having such a microstructure under predetermined conditions, the proportion of needle-shaped austenite that undergoes reverse transformation from nucleation sites within prior austenite grains can be increased. Furthermore, by applying appropriate treatment under predetermined cooling conditions, it is possible to obtain fine needle-shaped retained austenite with a relatively high concentration of alloying elements in the final obtained microstructure. More specifically, the martensite structure has substructures such as packets, blocks, and laths within the prior austenite grains, and therefore has many different interfaces within the grains compared to structures such as ferrite. Accordingly, by constructing the microstructure of the steel sheet before secondary heat treatment mainly with martensite, it is possible to increase the proportion of needle-shaped austenite that undergoes reverse transformation using various interfaces within the prior austenite grains as nucleation sites by heating such a microstructure to the two-phase region of ferrite and austenite during secondary heat treatment.

[0047] If the middle side of Equation 1 above, i.e., Mn / (Al+0.5Si), is 4.00 or less, Si and Al can be sufficiently segregated at the prior austenite grain boundaries in the martensitic structure before secondary heat treatment, thereby lowering the interfacial energy of the prior austenite grain boundaries. As a result, austenite can be preferentially nucleated from various interfaces within the prior austenite grains (such as packet and block boundaries) during secondary heat treatment, increasing the proportion of acicular austenite. On the other hand, if the middle side of Equation 1 above is greater than 4.00, Si and Al cannot be sufficiently segregated at the prior austenite grain boundaries in the martensitic structure before secondary heat treatment, making it impossible to preferentially nucleate austenite from various interfaces within the prior austenite grains during secondary heat treatment. As a result, the proportion of massive austenite becomes higher than that of acicular austenite, and it becomes impossible to obtain retained austenite with the desired X-ray diffraction peak characteristics in the final metal structure. Therefore, Mn / (Al+0.5Si) should be 4.00 or less, and may be, for example, 3.80 or less, 3.50 or less, 3.20 or less, 3.00 or less, 2.80 or less, 2.60 or less, or 2.40 or less.

[0048] On the other hand, before secondary heat treatment, the metal structure needs to be composed mainly of martensite, and in connection with this, the steel sheet needs to be heated to a high temperature of Ac3 or higher during primary heat treatment. However, when the steel sheet is heated at such a high temperature, coarse austenite is formed due to grain growth, and the metal structure obtained after subsequent cooling will naturally be composed mainly of coarse martensite. In such cases, even if the subsequent secondary heat treatment is performed appropriately, the retained austenite obtained will also be coarser. In the steel sheet according to the embodiment of the present invention, when the middle side of the above formula 1, i.e., Mn / (Al+0.5Si), is less than 1.00, the Ac3 point of the steel sheet rises, and this tendency becomes particularly pronounced, and it may become impossible to obtain retained austenite with the desired X-ray diffraction peak characteristics in the final metal structure. Alternatively, if Mn / (Al+0.5Si) is less than 1.00, Si and Al may excessively segregate at the prior austenite grain boundaries in the martensite structure before secondary heat treatment, causing the interfacial energy to drop too low. In this case, the amount of austenite that undergoes reverse transformation from the prior austenite grain boundaries during secondary heat treatment decreases significantly. As a result, the amount of austenite generated from the entire martensite structure becomes insufficient, reducing the area ratio of hard phases such as fresh martensite, tempered martensite, and bainite obtained after secondary heat treatment. Related to this, the area ratio of ferrite increases, and the desired strength may not be achieved. Therefore, in order to suppress austenite grain growth during primary heat treatment by controlling the Ac3 point to a relatively low value, thereby achieving refinement of the retained austenite ultimately obtained, and furthermore, to optimize the phase transformation to austenite and achieve the desired strength, Mn / (Al+0.5Si) needs to be controlled to 1.00 or higher. From the viewpoint of further refining the retained austenite, a larger value of Mn / (Al+0.5Si) is preferable. For example, Mn / (Al+0.5Si) may be 1.20 or higher, 1.40 or higher, 1.60 or higher, 1.80 or higher, 2.00 or higher, or 2.20 or higher.

[0049] [Metal structure] [Residual austenite: 5-35%] In the steel sheet according to the embodiment of the present invention, the metallic structure contains retained austenite: 5-35% by area. Retained austenite improves the uniform elongation of the steel sheet through the TRIP effect, where it transforms into a hard phase such as martensite by work-induced transformation during deformation of the steel sheet. To fully obtain this effect, the area ratio of retained austenite is set to 5% or more. From the viewpoint of further improving uniform elongation, a higher area ratio of retained austenite is preferable, for example, it may be 6% or more, 7% or more, 8% or more, 9% or more, 10% or more, 12% or more, 14% or more, or 15% or more. However, if the area ratio of retained austenite becomes too high, the area ratio of hard phases such as fresh martensite decreases, and therefore the desired strength may not be achieved. Therefore, the area ratio of retained austenite is set to 35% or less. The area ratio of retained austenite may be 30% or less, 28% or less, 26% or less, 24% or less, 22% or less, or 20% or less.

[0050] [Soft ferrite: 20-70%] Ferrite has a soft structure and is therefore easily deformable, contributing to improved ductility. In this invention, "soft ferrite" means ferrite that does not contain retained austenite within its grains. To fully obtain the effect of improving ductility, the area ratio of soft ferrite should be 20% or more. From the viewpoint of further improving ductility, a higher area ratio of soft ferrite is preferable, for example, it may be 25% or more, 30% or more, 35% or more, 40% or more, 45% or more, or 50% or more. On the other hand, as mentioned above, soft ferrite does not contain retained austenite within its grains, so its strength is lower. Therefore, if soft ferrite is included in excess, it may not be possible to achieve the desired strength in the steel sheet. For this reason, the area ratio of soft ferrite should be 70% or less. For example, the area ratio of soft ferrite may be 68% or less, 65% or less, 62% or less, 60% or less, 58% or less, or 55% or less.

[0051] [Fresh martensite: 8-35%] Fresh martensite is a rigid structure with a high dislocation density, and therefore contributes to improved tensile strength. To fully obtain this effect, the area ratio of fresh martensite should be 8% or more. From the viewpoint of further improving strength, a higher area ratio of fresh martensite is preferable, for example, 10% or more, 11% or more, 12% or more, 13% or more, 14% or more, 15% or more, 16% or more, 17% or more, 18% or more, 20% or more, or 22% or more. On the other hand, if the area ratio of fresh martensite is excessively high, ductility may decrease. Therefore, the area ratio of fresh martensite should be 35% or less. For example, the area ratio of fresh martensite may be 32% or less, 30% or less, 28% or less, or 25% or less. In this invention, "fresh martensite" means as-quenched martensite and does not include tempered martensite.

[0052] [Total of tempered martensite and bainite: 0-60%] In the steel sheet according to the embodiment of the present invention, the metal structure may include tempered martensite and / or bainite. Tempered martensite is a structure that improves the tensile strength of the steel sheet without impairing its impact resistance. Bainite is also a relatively hard structure and can similarly contribute to improving strength. Here, bainite includes granular bainite consisting of fine BCC crystals and coarse iron-based carbides, upper bainite consisting of lath-like BCC crystals and coarse iron-based carbides, and lower bainite consisting of plate-like BCC crystals and fine iron-based carbides arranged parallel to each other inside. The total area ratio of tempered martensite and bainite may be 0%, or for example, 3% or more, 5% or more, 10% or more, 15% or more, 20% or more, 25% or more, 30% or more, or 35% or more. On the other hand, excessive inclusion of tempered martensite and / or bainite may reduce ductility. Therefore, the total area ratio of tempered martensite and bainite should be 60% or less, and may be, for example, 55% or less, 50% or less, 45% or less, or 40% or less.

[0053] [Remaining tissue] As described above, the present invention aims to provide a steel sheet that, despite its high strength, exhibits improved ductility and work hardening ability in the later stages of deformation. This objective is achieved by controlling the microstructure of a steel sheet having a predetermined chemical composition so that its Vickers hardness is 300 Hv or higher, by configuring the microstructure to include predetermined amounts of retained austenite, soft ferrite, and fresh martensite, and by appropriately controlling the morphology of the retained austenite observed by X-ray diffraction. Therefore, in the microstructure of the steel sheet, the remaining microstructure other than the retained austenite, soft ferrite, fresh martensite, tempered martensite, and bainite described above is not particularly limited as long as it has a Vickers hardness of 300 Hv or higher, and is clearly not an essential technical feature for achieving the objective of the present invention. Accordingly, in the steel sheet according to the embodiment of the present invention, the remaining microstructure other than retained austenite, soft ferrite, fresh martensite, tempered martensite, and bainite can be appropriately selected within a range that results in a Vickers hardness of 300 Hv or higher for the steel sheet. The following provides a more detailed explanation of the tissues that can constitute the remaining tissue.

[0054] The remaining microstructure other than retained austenite, soft ferrite, fresh martensite, tempered martensite, and bainite may have an area ratio of 0%. If the remaining microstructure is present, it may contain pearlite and cementite. The remaining microstructure may also contain "hard ferrite" that encapsulates retained austenite within its grains. Because hard ferrite encapsulates retained austenite within its grains, it can contribute to improved strength. Furthermore, compared to cases where retained austenite is present at the ferrite grain boundaries, hard ferrite is less prone to interfacial delamination between ferrite and retained austenite, thus providing good moldability. From the viewpoint of ensuring the effects based on retained austenite, soft ferrite, fresh martensite, tempered martensite, and bainite, the area ratio of the remaining microstructure is preferably 20% or less in total, and may be, for example, 15% or less, 12% or less, 10% or less, 8% or less, 6% or less, 4% or less, 3% or less, or 2% or less. On the other hand, the area ratio of the remaining tissue may be 0.5% or more, or 1% or more. Also, the area ratio of hard ferrite may be, for example, 0-20%, and may be 0-15%, 0-12%, 0-10%, 0-8%, 0-6%, 0-4%, 0-3%, or 0-2%. The area ratio of hard ferrite may be 0.5% or more, or 1% or more.

[0055] [Identification of metallographic structure and calculation of area ratio] The identification of the metal structure and calculation of the area ratio are performed as follows: First, a sample is taken from the plate thickness cross section perpendicular to the plate surface, and this observation surface is polished and etched with nital. Next, in one or more observation fields in the range of 1 / 8 to 3 / 8 thickness centered at a position 1 / 4 thickness from the surface on the observation surface, a total of 2.0 × 10⁻⁶ -9 m 2The above areas are observed using a field emission scanning electron microscope (FE-SEM). The area percentages of ferrite, bainite, tempered martensite, fresh martensite, pearlite, and cementite are then measured. Here, regions with a substructure within the grain and where carbides precipitate with multiple variants are identified as tempered martensite. Regions where cementite precipitates in a lamellar pattern are identified as pearlite, and spherical cementite existing separately from the pearlite is identified as cementite. Regions with low brightness and no visible substructure are identified as ferrite (soft ferrite and hard ferrite). Regions with high brightness and where the substructure has not been revealed by etching are identified as fresh martensite and retained austenite. The remainder is identified as bainite. The area percentage of each structure is calculated using the point counting method. The area ratio of fresh martensite is determined by subtracting the area ratio of retained austenite, which is determined by the X-ray diffraction method described later, from the total area ratio of fresh martensite and retained austenite. The area ratios of hard ferrite and soft ferrite are determined based on the measured area ratio of ferrite using the method described later.

[0056] The area fraction of retained austenite is measured by X-ray diffraction. Specifically, a surface parallel to the surface of the steel plate is polished to a mirror finish in the range of 1 / 8 to 3 / 8 thickness, centered at a position 1 / 4 thickness from the plate surface. The integrated intensity ratio of the X-ray diffraction peaks at the (200), (211), and (200), (220), and (311) planes of the bcc phase and fcc phase is obtained using MoKα1 as the characteristic X-ray with a step width of 0.01° on the mirror-finished sample. The structural fraction of retained austenite is calculated from the integral intensity ratio of the X-ray diffraction peaks at the (200), (211), and (200), (220), and (311) planes of the fcc phase, and this is determined as the area fraction of retained austenite.

[0057] This paper describes a method for separating soft ferrite, which does not contain retained austenite, from hard ferrite, which does contain retained austenite. First, the crystal grains are observed using FE-SEM, and high-resolution crystal orientation analysis is performed by EBSD. Specifically, a sample is taken with the thickness cross section perpendicular to the plate surface as the observation surface, and the observation surface is polished to a mirror finish. Next, a total of 2.0 × 10⁻¹⁶ thickness is observed in one or more observation fields in the range from 1 / 8 thickness to 3 / 8 thickness, centered at a position 1 / 4 thickness from the surface. -9 m 2 Crystal structure analysis is performed on the above area using the EBSD method. Next, using the data obtained from BCC iron, a grain boundary map of ferrite is drawn, with boundaries resulting in a crystal orientation difference of 15° or more being defined as grain boundaries. Then, using the data obtained from FCC iron, a grain distribution map is drawn using only austenite grains with a major axis length of 0.1 μm or more to avoid measurement errors, and this is overlaid on the grain boundary map of ferrite grains. A ferrite grain is judged to be "hard ferrite" if it contains one or more austenite grains that are completely incorporated inside it. A ferrite grain is judged to be "soft ferrite" if it is not adjacent to austenite grains or is adjacent to austenite grains only at the boundary with other grains.

[0058] [Ratio of integrated intensity / peak height of X-ray diffraction peaks in the (220) plane of the fcc phase: 0.100 deg. or less, and ratio of integrated intensity / peak height of X-ray diffraction peaks in the (311) plane of the fcc phase: 0.120 deg. or less] In the steel sheet according to the embodiment of the present invention, the ratio of the integrated intensity / peak height of the X-ray diffraction peak at the (220) plane of the fcc phase at a position 1 / 4 of the sheet thickness from the surface of the steel sheet (hereinafter also simply referred to as "the ratio of the integrated intensity / peak height of the X-ray diffraction peak at the (220) plane of the fcc phase") is controlled to 0.100 deg. or less, and the ratio of the integrated intensity / peak height of the X-ray diffraction peak at the (311) plane of the fcc phase at a position 1 / 4 of the sheet thickness from the surface of the steel sheet (hereinafter also simply referred to as "the ratio of the integrated intensity / peak height of the X-ray diffraction peak at the (311) plane of the fcc phase") is controlled to 0.120 deg. or less.

[0059] By controlling the ratio of integrated intensity to peak height of the X-ray diffraction peak on the (220) plane of the fcc phase and the ratio of integrated intensity to peak height of the X-ray diffraction peak on the (311) plane of the fcc phase to the above range, that is, by controlling the X-ray diffraction peaks on the (220) and (311) planes of the fcc phase to a relatively sharp shape, it is possible to sufficiently form fine needle-shaped retained austenite with relatively high concentrations of alloying elements such as C and Mn. As a result, the ductility of the steel sheet can be further improved, and the work hardening rate in the high strain region can be significantly increased, thereby making it possible to achieve a high work hardening rate even in conditions where some strain has been introduced, such as in the later stages of deformation during press forming.

[0060] For example, if the ratio of integrated intensity to peak height of the X-ray diffraction peak on the (220) plane of the fcc phase is greater than 0.100 deg. and / or the ratio of integrated intensity to peak height of the X-ray diffraction peak on the (311) plane of the fcc phase is greater than 0.120 deg., the peak heights of these diffraction peaks are low, resulting in a relatively large proportion of retained austenite having a massive rather than needle-like form. Consequently, the concentrations of alloying elements such as C and Mn in the retained austenite also decrease. As a result, the ductility of the steel sheet cannot be further improved, and the work hardening ability in the later stages of deformation cannot be sufficiently improved. For example, even if the ratio of integrated intensity to peak height of the X-ray diffraction peak on the (220) plane of the fcc phase is controlled to be 0.100 deg. or less, if the ratio of integrated intensity to peak height of the X-ray diffraction peak on the (311) plane of the fcc phase is not controlled to be 0.120 deg. or less, fine needle-like retained austenite with relatively high concentrations of alloying elements may not be sufficiently formed. Even in such cases, it may not be possible to fully obtain the effect of further improving ductility and improving work hardening ability in the later stages of deformation. Therefore, in the steel sheet according to the embodiment of the present invention, in order to reliably obtain the above effects, it is important to control the ratio of integrated intensity / peak height of the X-ray diffraction peak on the (220) plane of the fcc phase to 0.100 deg. or less, and to control the ratio of integrated intensity / peak height of the X-ray diffraction peak on the (311) plane of the fcc phase to 0.120 deg. or less. Although it is not intended to be bound by any particular theory, it is thought that when the amount of fine needle-shaped retained austenite with relatively high concentrations of C and Mn, which contribute to improved ductility, increases, the peaks in X-ray diffraction that were spread out with the superposition of retained austenite having multiple compositions (multiple compositions from low to high alloying element concentrations) converge to retained austenite having a single composition (composition with medium alloying element concentrations), and the peaks become sharper. In embodiments of the present invention, such peak sharpening can be expressed using an index called the "ratio of integrated intensity of X-ray diffraction peak to peak height".

[0061] To further enhance the above effects, a lower ratio of integrated intensity to peak height for the X-ray diffraction peaks on each plane is preferable. Therefore, the ratio of integrated intensity to peak height for the X-ray diffraction peaks on the (220) plane of the fcc phase may be, for example, 0.098 deg. or less, 0.096 deg. or less, 0.094 deg. or less, 0.092 deg. or less, or 0.090 deg. or less. Similarly, the ratio of integrated intensity to peak height for the X-ray diffraction peaks on the (311) plane of the fcc phase may be, for example, 0.118 deg. or less, 0.116 deg. or less, 0.114 deg. or less, 0.112 deg. or less, 0.110 deg. or less, or 0.108 deg. or less. The lower limit is not particularly limited, but for example, the ratio of integrated intensity to peak height of the X-ray diffraction peak on the (220) plane of the fcc phase may be 0.070 deg. or higher, 0.075 deg. or higher, or 0.080 deg. or higher. Similarly, the ratio of integrated intensity to peak height of the X-ray diffraction peak on the (311) plane of the fcc phase may be 0.090 deg. or higher, 0.095 deg. or higher, or 0.100 deg. or higher.

[0062] [Method for determining the ratio of integrated intensity to peak height of X-ray diffraction peaks in the (220) plane of the fcc phase and the ratio of integrated intensity to peak height of X-ray diffraction peaks in the (311) plane of the fcc phase] The ratio of the integrated intensity to peak height of the X-ray diffraction peak at the (220) plane of the fcc phase at a position 1 / 4 of the plate thickness from the steel plate surface, and the ratio of the integrated intensity to peak height of the X-ray diffraction peak at the (311) plane of the fcc phase at a position 1 / 4 of the plate thickness from the steel plate surface, are determined by obtaining the integrated intensity (cps·deg.) and their peak heights (cps) of the X-ray diffraction peaks at the (220) and (311) planes of the fcc phase during the measurement of the area fraction of retained austenite as described in the section on [Identification of Metal Structure and Calculation of Area Ratio], and then calculating their ratios. In measuring the X-ray diffraction peak, background is removed from the X-ray diffraction peak using software such as "Quantitative Analysis of Retained Austenite" from Rigaku Corporation. Then, the integrated intensity I of the X-ray diffraction peak (sum of Kα1 and Kα2 lines) is determined. Since the integrated intensity ratio of Kα1 and Kα2 lines is observed to be 2:1, the integrated intensity of the Kα1 line X-ray diffraction peak can be calculated as I × 2 / 3. The peak height is the maximum value of the X-ray diffraction peak after background removal. If the steel sheet has a plating layer, the plating layer is removed before measuring the X-ray diffraction peak. Specifically, for example, the plating layer can be removed from the steel sheet by dissolving it with an acidic aqueous solution containing an inhibitor that suppresses the dissolution of steel, such as a room-temperature acidic aqueous solution of 10% hydrochloric acid with 0.04% Ibit 710K (manufactured by Asahi Chemical Industry Co., Ltd.) added.

[0063] [Thickness of the soft layer: 1-100 μm] In a preferred embodiment of the present invention, the steel sheet includes a soft layer having a thickness of 1 to 100 μm in the thickness direction from the surface of the steel sheet, where the soft layer is defined as a region having a hardness of 80% or less of the hardness at a position 1 / 4 of the thickness from the surface of the steel sheet. Including such a soft layer in the surface layer of the steel sheet makes it possible to further improve the ductility of the steel sheet. In addition, although the reason is not entirely clear, research by the inventors has shown that carbon contained in the steel sheet is an element that promotes LME cracking. Therefore, by providing a soft layer with a reduced carbon concentration in the surface layer of a steel sheet where LME cracking occurs, it is possible to more significantly suppress or reduce the occurrence of LME cracking compared to simply reducing the Si content of the steel sheet.

[0064] From the viewpoint of further improving the ductility and LME resistance of the steel sheet, a thicker soft layer is preferable. Therefore, when the soft layer is defined as the region having a hardness of 80% or less of the hardness at a position 1 / 4 of the sheet thickness from the surface of the steel sheet, the thickness of the soft layer may be 5 μm or more, 10 μm or more, 20 μm or more, 30 μm or more, 40 μm or more, 50 μm or more, or 60 μm or more in the direction of the sheet thickness from the surface of the steel sheet. However, since the above effect saturates even if the soft layer is made thicker than 100 μm, in a preferred embodiment of the present invention, the thickness of the soft layer is 100 μm or less, and may be, for example, 90 μm or less, 80 μm or less, or 70 μm or less.

[0065] [Determination of the thickness of the soft layer] The thickness of the soft layer is determined by Vickers hardness in accordance with JIS Z 2244-1:2024 as follows: First, a test piece is cut from the steel plate and embedded in resin so that a cross-section perpendicular to the plate thickness can be observed. Then, the embedded surface is polished using silicon carbide sandpaper from #600 to #1500, and then finished to a mirror surface using a liquid in which diamond powder with a particle size of 1 to 6 μm is dispersed in a diluent such as alcohol or pure water. The cross-section embedded in this resin is used as the measurement surface. Using a microhardness analyzer, measurements are taken at a test force of 0.4903 N (i.e., HV 0.05) at intervals of at least three times the diagonal length of the indentation, in an oblique direction from the surface (for example, at a 10° gradient to the steel plate surface). After that, the diagonal length of the indentation is measured using a scanning electron microscope (SEM) and the Vickers hardness (Hv) is calculated. The thickness of the soft layer (the region having a hardness of 80% or less of the hardness at the 1 / 4 thickness position from the steel plate surface) is obtained in μm units by linear interpolation between each measurement value. However, the hardness at the 1 / 4 thickness position is measured at 5 points at 0.1 mm intervals, and the arithmetic mean of these 5 measurement values ​​is taken as the hardness at the 1 / 4 thickness position. If the obtained soft layer thickness is less than 1.0 times the diagonal length of the depression, the soft layer thickness is judged to be 0 μm.

[0066] [Vickers hardness: 300Hv or higher] The steel sheet according to the embodiment of the present invention has a Vickers hardness of 300 Hv or more. Preferably, the Vickers hardness is 310 Hv or more, 330 Hv or more, 350 Hv or more, 360 Hv or more, or 380 Hv or more. According to the embodiment of the present invention, by controlling the microstructure of a steel sheet having a predetermined chemical composition so that the Vickers hardness is 300 Hv or more, and by constructing the microstructure with a predetermined amount of retained austenite, and further appropriately controlling the morphology of the retained austenite observed by X-ray diffraction, it is possible to improve ductility and work hardening ability in the later stages of deformation despite high strength. The upper limit of the Vickers hardness is not particularly limited, but for example, the Vickers hardness may be 600 HV or less, 580 HV or less, 550 HV or less, or 520 HV or less. The Vickers hardness of a steel sheet refers to the hardness at the 1 / 4 thickness position calculated when determining the thickness of the soft layer.

[0067] [plate thickness] The steel plate according to the embodiment of the present invention is not particularly limited, but generally has a thickness of 0.6 to 8.0 mm. For example, the thickness may be 0.8 mm or more, 1.0 mm or more, 1.2 mm or more, 1.4 mm or more, 1.6 mm or more, or 2.0 mm or more, and / or 7.0 mm or less, 6.0 mm or less, 5.0 mm or less, 4.0 mm or less, 3.0 mm or less, or 2.5 mm or less.

[0068] [Plating layer] A steel sheet according to an embodiment of the present invention may have a zinc-containing plating layer on at least one surface, preferably both surfaces. The plating layer is not particularly limited, but may be, for example, a hot-dip galvanized layer (GI), an alloyed hot-dip galvanized layer (GA), or an electroplated galvanized layer (EG). Here, a zinc plating layer refers to a plating layer with a Zn content of more than 50%. These zinc plating layers may have any chemical composition known to those skilled in the art within the range of a Zn content of more than 50%, and may contain additive elements such as Al or Mg in addition to Zn. In addition to zinc plating layers, the plating layer may also be a hot-dip Al plating layer. For example, the chemical composition of the plating layer can be obtained by dissolving the plating layer in an acidic aqueous solution containing an inhibitor that suppresses the dissolution of steel, for example, an acidic aqueous solution at room temperature with 10% hydrochloric acid and 0.04% Ibit 710K (manufactured by Asahi Chemical Industry Co., Ltd.), and analyzing the resulting acidic aqueous solution by ICP (inductively coupled plasma) emission spectroscopy. Furthermore, the amount of the plating layer is not particularly limited and may be a general amount. The steel sheets according to the embodiments of the present invention are, of course, not limited to plated steel sheets, but also include unplated steel sheets. Even with unplated steel sheets, for example, when spot welding with a galvanized steel sheet, molten zinc from the galvanized steel sheet may penetrate into the unplated steel sheet, causing LME cracking. However, the steel sheets according to the embodiments of the present invention have improved LME resistance by limiting the Si content to 2.00 mass% or less, making them useful in the automotive field where high weldability is required. In addition, when a soft layer having a thickness of 1 to 100 μm in the thickness direction from the surface of the steel sheet is included as described above, it is possible to suppress or reduce the occurrence of LME cracking more significantly compared to when the Si content of the steel sheet is simply limited to 2.00 mass% or less.

[0069] As described above, the steel sheet according to the embodiment of the present invention controls the metal structure of a steel sheet having a predetermined chemical composition so that its Vickers hardness is 300 Hv or higher, and the metal structure is composed of a structure containing a predetermined amount of retained austenite. Furthermore, by appropriately controlling the morphology of the retained austenite observed by X-ray diffraction, it is possible to improve ductility and work hardening ability in the later stages of deformation, despite high strength. Therefore, the steel sheet according to the embodiment of the present invention is particularly useful for use in parts in technical fields that require a high level of both high strength and formability, and is especially useful for use in parts in the automotive field. In a preferred embodiment, an automotive part (including not only passenger cars but also parts used in large vehicles such as trucks) containing the steel sheet according to the embodiment of the present invention is provided. Examples of automotive parts include skeletal parts such as front pillars, center pillars, side sills, and cross members, as well as bumpers, and other structural and reinforcing parts that require strength. These parts only need to contain the steel sheet according to the embodiment of the present invention in at least a portion of them, and therefore at least a portion of these parts will satisfy the characteristics of the steel sheet described above. In forming processes such as press forming, the characteristics of the steel sheet do not change significantly before and after forming in areas of the steel sheet that do not come into direct contact with the mold, or in areas where the degree of processing is relatively low even if they come into direct contact with the mold.

[0070] [Mechanical properties] [Tensile strength (TS) and uniform elongation (uEL)] According to the steel sheet having the above chemical composition and steel structure, a high tensile strength, specifically a tensile strength (TS) of 980 MPa or higher, can be achieved. The tensile strength is preferably 1080 MPa or higher, 1180 MPa or higher, or 1200 MPa or higher. There is no particular upper limit, but for example, the tensile strength may be 1800 MPa or lower, 1600 MPa or lower, 1500 MPa or lower, or 1400 MPa or lower. Furthermore, according to the steel sheet according to the embodiment of the present invention, despite having such a very high tensile strength, an improved elongation, specifically a uniform elongation (uEL) of 12.0% or higher, can be achieved. For example, the uniform elongation may be 13.0% or higher, 14.0% or higher, 15.0% or higher, or 16.0% or higher. For example, when the tensile strength is 980 MPa or higher, a uniform elongation of 15.0% or higher can be achieved. Similarly, for example, when the tensile strength is 1180 MPa or higher, a uniform elongation of 12.0% or higher can be achieved. The upper limit of uniform elongation is not particularly limited, but for example, the uniform elongation may be 25.0% or less, 22.0% or less, or 20.0% or less. Furthermore, according to the steel sheet according to the embodiment of the present invention, ductility can also be evaluated by TS×uEL, and for example, it is possible to achieve a TS×uEL of 14140MPa·% or more, 15000MPa·% or more, 16000MPa·% or more, 17000MPa·% or more, or 18000MPa·% or more. The upper limit is not particularly limited, but for example, the TS×uEL may be 25000MPa·% or less or 24000MPa·% or less. Tensile strength and uniform elongation are determined by taking a JIS No. 5 test specimen from a direction (C direction) where the longitudinal direction of the test specimen is preferably parallel to the direction perpendicular to the rolling direction of the steel sheet, and performing a tensile test in accordance with JIS Z 2241:2022. Here, uniform elongation refers to "plastic elongation at maximum test force (%)Ag" as defined in JIS Z 2241:2022. If the rolling direction of the steel sheet cannot be determined, a JIS No. 5 test specimen may be taken from any direction within the surface of the steel sheet. If it is difficult to take a JIS No. 5 test specimen, a JIS No. 13B test specimen may be used, or a small test specimen with a similar shape to a JIS No. 13B test specimen may be used.

[0071] <Method of manufacturing steel plates> Next, preferred manufacturing methods for steel sheets according to embodiments of the present invention will be described. The following description is intended to illustrate characteristic methods for manufacturing steel sheets according to embodiments of the present invention, and is not intended to limit the steel sheets to those manufactured by the manufacturing methods described below.

[0072] A method for manufacturing steel sheets according to an embodiment of the present invention is: The heating process involves heating a slab having the chemical composition described above in relation to a steel plate to a temperature of 1230-1350°C and holding it for 1.5-3.0 hours, satisfying the following formula 2, where w and D in formula 2 are the heating steps represented by the following formulas 3 and 4.

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[0073] [Soaking process] Optionally, a soaking treatment may be performed before the heating process, in which the slab is heated to a temperature of 1230-1350°C and held for 1.5-40 hours. Performing such a soaking treatment before the heating process makes it possible to further improve the ductility of the final steel sheet compared to cases where the soaking treatment is not performed.

[0074] [Heating process] First, a slab having the chemical composition described above in relation to the steel sheet is heated. From the viewpoint of productivity, the slab used is preferably cast by a continuous casting method, but it may also be manufactured by an ingot-making method or a thin-slab casting method. The slab used contains a relatively large amount of alloying elements such as Mn in order to improve the hardenability of the steel sheet. Mn is an element that tends to segregate in streaks in the steel sheet, and more specifically, Mn-enriched regions such as central segregation and microsegregation are formed during casting. When such Mn-enriched regions exist, a relatively large amount of band-shaped retained austenite is formed in the final metal structure, and the amount of Mn distributed in the retained austenite becomes non-uniform, making it impossible to obtain retained austenite with the desired X-ray diffraction peak characteristics. Therefore, in this manufacturing method, the slab is heated to a temperature of 1230 to 1350°C and held for 1.5 to 3.0 hours, and then subjected to a heating step that satisfies the following formula 2, thereby allowing Mn to diffuse sufficiently in the slab, eliminating central and microsegregation of Mn and making the Mn concentration distribution in the slab uniform. As a result, the amount of Mn distributed in the retained austenite in the final metal structure can be made uniform. Here, w and D in equation 2 below are expressed in equations 3 and 4 below, more specifically, equation 3 below is an equation relating to the Mn segregation width, and equation 4 below is an equation relating to the diffusion of Mn.

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[0075] If the heating temperature is less than 1230°C and / or the holding time is less than 1.5 hours, Mn may not diffuse sufficiently in the slab, resulting in a non-uniform Mn concentration distribution in the slab. In such cases, even if subsequent heat treatment is performed appropriately, the distribution of retained austenite after the secondary heat treatment process will be band-like, and the amount of Mn distributed in the retained austenite will also be non-uniform, making it impossible to obtain retained austenite with the desired X-ray diffraction peak characteristics. As a result, the desired ductility and / or work hardening ability in the later stages of deformation cannot be obtained. Preferably, the heating temperature is 1250°C or higher and the holding time is 1.6 hours or higher. On the other hand, if the heating temperature exceeds 1350°C and / or the holding time exceeds 3.0 hours, the load on the heating equipment becomes excessive, and productivity decreases. Preferably, the heating temperature is 1330°C or lower and the holding time is 2.9 hours or lower. On the other hand, if the value of the left side of Equation 2 is greater than 0.70, the Mn concentration distribution in the slab becomes non-uniform, and it becomes impossible to obtain retained austenite with the desired X-ray diffraction peak characteristics in the final metal structure. Preferably, the value of the left side of Equation 2 is 0.65 or less.

[0076] [Hot rolling process] [Rough rolling] In this manufacturing method, for example, a heated slab may be subjected to rough rolling before finish rolling to adjust the plate thickness. The conditions for rough rolling are not particularly limited, as long as the desired sheet bar dimensions are ensured.

[0077] [Finishing rolling and winding] The heated slab, or slab that has been roughly rolled as needed, is then subjected to finish rolling. As the slabs used as described above contain a relatively large amount of alloying elements such as Mn, it is necessary to increase the rolling load during hot rolling. For this reason, it is preferable to perform hot rolling at a high temperature. In particular, the end temperature of finish rolling is important in terms of controlling the metal structure of the steel sheet. If the end temperature of finish rolling is low, the metal structure may become non-uniform, and the formability may decrease. For this reason, the end temperature of finish rolling should be 800°C or higher. On the other hand, in order to suppress the coarsening of austenite, the end temperature of finish rolling should be 1250°C or lower. Next, the hot-rolled steel sheet that has been finish-rolled is wound at a winding temperature of 500 to 700°C. By setting the winding temperature to 500 to 700°C, the growth of oxide scale can be suppressed.

[0078] [Pickling process] Next, the obtained hot-rolled steel sheet is pickled to remove the oxide scale formed on its surface. The pickling can be carried out under conditions suitable for removing the oxide scale, and may be done once or in multiple steps to ensure complete removal.

[0079] [Cold rolling process] Pickled hot-rolled steel sheets are cold-rolled in the cold-rolling process with a reduction ratio of 20-90%. By setting the cold-rolling reduction ratio to 20% or more, the shape of the cold-rolled steel sheet can be kept flat, suppressing a decrease in ductility in the final product. On the other hand, by setting the cold-rolling reduction ratio to 90% or less, it is possible to prevent excessive rolling loads that would make rolling difficult. The number of rolling passes and the reduction ratio for each pass are not particularly limited; they should be set appropriately so that the overall cold-rolling reduction ratio falls within the above range.

[0080] [Primary heat treatment process] The obtained cold-rolled steel sheet is heated in the following primary heat treatment step, held at a maximum heating temperature of Ac3 to 1000°C for 10 to 1000 seconds, and then cooled to a cooling stop temperature of 300°C or lower by controlling the average cooling rate in the temperature range of 500 to 700°C to 10°C / second or higher. In this manufacturing method, the Ac3 point (°C) is calculated based on the following formula 5. In the following formula 5, the mass % of the element is substituted for the element symbol. For elements that are not present, 0 mass % is substituted. Ac3(°C) = 910 - 203.0 × C 0.5 +44.7×Si-30.0×Mn+700.0×P-20.0×Cu-15.2×Ni-11.0×Cr+31.5×Mo+400.0×Ti+104.0×V+120.0×Al...Formula 5 By maintaining a temperature above the Ac3 point for a sufficient amount of time, austenitization is promoted, and by subsequent rapid cooling to a temperature below 300°C, it is possible to reliably establish a martensite-dominant microstructure in the steel sheet after cooling, such as full martensite. Here, a martensite-dominant microstructure refers to a microstructure containing 90% or more martensite, and full martensite refers to a microstructure containing 100% martensite. A martensite microstructure has substructures such as packets, blocks, and laths within the prior austenite grains, and therefore has many different interfaces within the grains compared to microstructures such as ferrite. For this reason, by mainly constructing the microstructure of the steel sheet before the secondary heat treatment process with martensite, it is possible to increase the proportion of needle-shaped austenite that undergoes reverse transformation using the various interfaces within the prior austenite grains as nucleation sites when such a microstructure is heated to the two-phase region of ferrite and austenite during the secondary heat treatment. As a result, in the metal structure after subsequent secondary heat treatment, it becomes possible to sufficiently form retained austenite with the desired X-ray diffraction peak characteristics, i.e., fine, needle-shaped retained austenite with relatively high concentrations of alloying elements such as C and Mn.

[0081] If the primary heat treatment process is omitted, it is naturally impossible to construct a martensite-dominant microstructure in the steel sheet before the secondary heat treatment process. Furthermore, even if the primary heat treatment process is performed, if the maximum heating temperature in the primary heat treatment process is below the Ac3 point or the holding time is less than 10 seconds, austenitization will be insufficient, and even after subsequent cooling, it will not be possible to construct a martensite-dominant microstructure in the steel sheet. In other words, it will not be possible to achieve a martensite area ratio of 90% or more. As a result, it will not be possible to obtain retained austenite with the desired X-ray diffraction peak characteristics in the final microstructure. On the other hand, heating and holding at higher temperatures and for longer periods will coarseen the austenite due to grain growth and further reduce productivity, therefore, the maximum heating temperature in the primary heat treatment process should be 1000°C or less, and the holding time should be 1000 seconds or less.

[0082] Furthermore, if the average cooling rate in the 500-700°C temperature range during the primary heat treatment process is less than 10°C / second, or if the cooling stop temperature exceeds 300°C, ferrite will form during cooling, making it impossible to achieve a martensite area ratio of 90% or more in the microstructure of the steel sheet. As a result, the area ratio of ferrite in the final microstructure will be high, making it impossible to achieve the desired strength, and / or to obtain sufficient retained austenite with the desired X-ray diffraction peak characteristics. Therefore, the average cooling rate must be 10°C / second or higher, with an upper limit of preferably 300°C / second or 150°C / second. On the other hand, the lower limit of the cooling stop temperature is not particularly limited and may be, for example, room temperature (about 20°C).

[0083] [Secondary heat treatment process] The cold-rolled steel sheet after primary heat treatment is heated in the subsequent secondary heat treatment process, held at a maximum heating temperature of (Ac1+25)~(Ac3-10)°C for 5~1000 seconds, then cooled to a temperature below 50°C, and then reheated to a temperature range of 300~450°C, where it is held for 100~500 seconds. In this manufacturing method, the Ac1 point (°C) is calculated based on the following formula 6. In formula 6 below, the mass % of the element is substituted for the element symbol. For elements that are not present, 0 mass % is substituted. Ac1(℃)=727-32.7×C+14.9×Si+2.0×Mn+17.8×Cr+25.6×Mo...Formula 6 First, by heating the cold-rolled steel sheet after primary heat treatment to a maximum heating temperature of (Ac1+25) to (Ac3-10)°C, corresponding to the two-phase region of ferrite and austenite, and holding it for 5 to 1000 seconds, it is possible to increase the proportion of needle-shaped austenites that undergo reverse transformation using various interfaces within the prior austenite grains as nucleation sites. During heating, it is preferable that the average heating rate from (Ac1-30)°C to the maximum heating temperature be 0.10 to 100.00°C / second. If the average heating rate from (Ac1-30)°C to the maximum heating temperature is less than 0.10°C / second, productivity will decrease. On the other hand, if the average heating rate exceeds 100.00°C / second, the number of austenite nucleations may decrease. More preferably, the average heating rate is 50.00°C / second or less, or 10.00°C / second or less. Next, by cooling the steel plate to a temperature below 50°C, relatively unstable austenite with low alloy concentrations of C and Mn is transformed into martensite, leaving only relatively stable austenite with high alloy concentrations. As a result, in the final resulting microstructure, the relatively stable austenite with high alloy concentrations remains as retained austenite, making it possible to form fine, needle-shaped retained austenite with relatively high concentrations of alloying elements, i.e., retained austenite with the desired X-ray diffraction peak characteristics.

[0084] If the maximum heating temperature in the secondary heat treatment step is less than (Ac1+25)°C or the holding time is less than 5 seconds, cementite in the steel may remain undissolved, resulting in insufficient austenite formation. This can lead to a decrease in the area ratio of retained austenite obtained after the secondary heat treatment, making it impossible to obtain the desired ductility. On the other hand, if the maximum heating temperature is higher than (Ac3-10)°C, austenitization proceeds excessively, and the lath-like structure of martensite obtained in the primary heat treatment step is not retained. In this case, the reverse transformation from the prior austenite grains cannot proceed sufficiently, and needle-shaped retained austenite cannot be sufficiently formed. As a result, retained austenite with the desired X-ray diffraction peak characteristics cannot be formed, and the desired ductility and / or work hardening ability in the later stages of deformation cannot be obtained. Preferably, the maximum heating temperature in the secondary heat treatment step is (Ac1+30) to (Ac3-15)°C.

[0085] In a preferred embodiment of this manufacturing method, the atmosphere during which the cold-rolled steel sheet is held at a temperature range of (Ac1+25) to (Ac3-10)°C for 5 to 1000 seconds is controlled to satisfy the following formula 7 and to have an H2 concentration of 0.1 to 20.0 volume percent.

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[0086] The middle side of equation 7 above, i.e., log(pH2O / pH2), is also called the oxygen potential, and the larger this value, the more the decarburization of the surface layer of the steel sheet can be advanced. In the secondary heat treatment process, by controlling the oxygen potential in the temperature range of (Ac1+25)~(Ac3-10)℃ to satisfy equation 7 above and the H2 concentration to 0.1~20.0 volume%, and further controlling the holding time in this temperature range to within the range of 5~1000 seconds, it is possible to appropriately decarburize the surface layer of the steel sheet. As a result, when the region having a hardness of 80% or less of the hardness at a position 1 / 4 of the thickness from the surface is defined as the soft layer in the final steel sheet, it is possible to ensure that the sheet includes a soft layer with a thickness of 1~100 μm in the thickness direction from the surface, thereby further improving the ductility of the steel sheet and more significantly suppressing or reducing the occurrence of LME cracks. From the viewpoint of further improving the ductility and LME resistance of the steel sheet, a thicker soft layer is preferable, and therefore a larger value of log(pH2O / pH2) is preferable. For example, log(pH2O / pH2) should be -1.8 or higher, but may also be -1.5 or higher, or -1.2 or higher. On the other hand, if the oxygen potential in the temperature range of (Ac1+25) to (Ac3-10)°C becomes excessively high, decarburization may proceed too much, which can lead to excessive softening of the surface layer of the steel sheet and a decrease in the strength of the steel sheet. For this reason, it is preferable to keep the oxygen potential in the temperature range of (Ac1+25) to (Ac3-10)°C at -0.4 or lower, and may also be -0.6 or lower.

[0087] If the H2 concentration in the above atmosphere is less than 0.1 volume%, an oxide film may form on the surface of the steel sheet, which may reduce the chemical conversion treatment properties of the steel sheet obtained after the secondary heat treatment process. Related to this, the adhesion of the plating when applying plating treatments such as hot-dip galvanizing may decrease. Therefore, the H2 concentration in the above atmosphere should be 0.1 volume% or more, preferably 1.0 volume% or more, and more preferably 2.0 volume% or more. On the other hand, if the H2 concentration in the above atmosphere becomes too high, the operational risks will naturally increase. For this reason, the H2 concentration should be 20.0 volume% or less, for example, 10.0 volume% or less or 5.0 volume% or less.

[0088] Furthermore, if the cooling stop temperature after heating to the maximum heating temperature is 50°C or higher, it becomes impossible to sufficiently transform relatively unstable austenite with low alloy concentrations of C and Mn into martensitic material. As a result, in the final resulting metal structure, it is not possible to form fine, needle-shaped retained austenite with relatively high concentrations of alloying elements, i.e., retained austenite having the desired X-ray diffraction peak characteristics, and the desired ductility and / or work hardening ability in the later stages of deformation cannot be obtained. The lower limit of the cooling stop temperature is not particularly limited and may be, for example, room temperature (approximately 20°C). By setting the lower limit of the cooling stop temperature to 5°C, the area ratio of fresh martensite can be controlled to 12% or higher. During cooling, it is preferable that the average cooling rate from the maximum heating temperature to (Ac1-150)°C is 2.0°C / second or higher. If the average cooling rate from the maximum heating temperature to (Ac1-150)°C is less than 2.0°C / second, a relatively large amount of soft ferrite may be formed, and the strength may decrease somewhat. Therefore, in order to ensure that the desired intensity is achieved, the average cooling rate is preferably 2.0°C / second or higher, and more preferably 2.5°C / second or higher. There is no particular upper limit, but for example, the average cooling rate from the maximum heating temperature to (Ac1-150)°C may be 20.0°C / second or lower, or 10.0°C / second or lower.

[0089] Finally, the cold-rolled steel sheet is reheated to a temperature range of 300-450°C and held at that temperature for 100-500 seconds. This allows the area ratio of fresh martensite and other elements to be controlled within the desired range, making it possible to achieve a good balance of strength and ductility. More specifically, by controlling the reheating temperature to 300-450°C, the tempering of martensite can be appropriately controlled, and the excessive formation of fresh martensite can be suppressed. In connection with this, the amount of retained austenite formed can also be controlled within the desired range. If the reheating temperature is below 300°C, the tempering of martensite cannot be promoted, the area ratio of fresh martensite may increase, and the ductility may decrease. Alternatively, if the tempering of martensite cannot be promoted, carbon enrichment into the untransformed austenite may not progress, and the desired area ratio of retained austenite may not be achieved in the final metal structure. On the other hand, if the reheating temperature exceeds 450°C, excessive tempering of the martensite may prevent the final steel sheet from achieving the desired strength, or / or a relatively large amount of carbides may be formed, preventing the achievement of the desired area ratio of retained austenite.

[0090] If the holding time in the 300-450°C temperature range is less than 100 seconds, bainite will not grow sufficiently, and carbon enrichment into the un-transformed austenite will not proceed. As a result, the stabilization of retained austenite will not proceed sufficiently, and a relatively large amount of fresh martensite may be produced, or / or retained austenite may not be produced sufficiently, which may reduce ductility. On the other hand, if the holding time in the 300-450°C temperature range is more than 500 seconds, the effect will saturate and productivity will decrease. Preferably, the holding time in the 300-450°C temperature range is 150-450 seconds.

[0091] [Plating process] When manufacturing plated steel sheets, they may be immersed in a plating bath (approximately 460°C in the case of a Zn bath) during the secondary heat treatment process, for example, while they are being heated to the two-phase region and then cooled to below 50°C. Alternatively, they may be heated after the secondary heat treatment process, for example, after being cooled to room temperature, and then immersed in a plating bath (approximately 460°C in the case of a Zn bath). When manufacturing alloyed hot-dip galvanized steel (GA), the steel may be reheated and alloyed after immersion in the plating bath. The alloying temperature may be in the range of 460 to 600°C. Alternatively, hot-dip galvanized steel (GI) without alloying may also be used. For example, the steel sheet after the secondary heat treatment process may be electro-galvanized (EG). These zinc plating processes may be carried out according to any suitable method known to those skilled in the art. Similarly, these zinc plating processes may have any composition known to those skilled in the art, and may contain additive elements such as Al and Mg in addition to Zn. Furthermore, the amount of zinc plating applied is not particularly limited and may be a general amount.

[0092] According to the steel sheet manufactured by the above manufacturing method, by limiting the Si content to 2.00 mass% or less while increasing the Al content to 0.200 mass% or more, and further controlling the Vickers hardness of the steel sheet to 300 Hv or more, while configuring the steel structure of the steel sheet to contain retained austenite: 5-35% by area%, it is possible to improve the ductility of the steel sheet while achieving high strength, for example, a tensile strength of 980 MPa or more. Furthermore, by using a chemical composition satisfying 1.00 ≤ Mn / (Al + 0.5Si) ≤ 4.00, a controlled heating process, and a specific combination of primary and secondary heat treatment processes, retained austenite can be obtained in which the ratio of integrated intensity to peak height of the X-ray diffraction peak on the (220) plane of the fcc phase is controlled to 0.100 deg. or less, and the ratio of integrated intensity to peak height of the X-ray diffraction peak on the (311) plane of the fcc phase is controlled to 0.120 deg. or less, i.e., fine needle-shaped retained austenite with relatively high concentrations of alloying elements such as C and Mn. As a result, the ductility of the steel sheet can be further improved, and the work hardening rate in the high-strain region can be significantly increased, thereby enabling a high work hardening rate even in conditions where some strain has been introduced, such as in the later stages of deformation during press forming. Therefore, steel sheets manufactured by the above manufacturing method are particularly useful in the automotive sector, where a high level of both high strength and formability is required.

[0093] The present invention will be described in more detail below with reference to examples, but the present invention is not limited in any way to these examples. [Examples]

[0094] In the following examples, steel sheets according to the embodiment of the present invention were manufactured under various conditions, and the tensile strength (TS), uniform elongation (uEL), and work hardening ability in the later stages of deformation of the obtained steel sheets were investigated.

[0095] First, molten steel was cast using a continuous casting method to form slabs with various chemical compositions (molten steel analysis values) as shown in Table 1. Equation 1 in Table 1 represents the value of "Mn / (Al+0.5Si)". Next, these slabs were arbitrarily soaked under the conditions shown in Table 2, and then heated and held under the same conditions as shown in Table 2. Next, the slabs were hot-rolled and then coiled at a temperature of 550-650°C. Hot rolling was performed in rough rolling and finish rolling stages, with the finish rolling completion temperature being 900-1100°C. Next, the obtained hot-rolled steel sheets were pickled and cold-rolled with a reduction ratio of 50-80%. The thickness of the cold-rolled sheets was 1.6 mm in all cases. The obtained cold-rolled steel sheets were heated and held under the conditions shown in Table 2 in the primary heat treatment stage, and then cooled. Next, the cold-rolled steel sheets after primary heat treatment were heated and held under the conditions shown in Table 2 during the secondary heat treatment process, and then cooled to the cooling stop temperature shown in Table 2. In connection with this heating and holding, the atmosphere during which the cold-rolled steel sheets were held in the temperature range of (Ac1+25) to (Ac3-10)°C was controlled to the oxygen potential and H2 concentration (volume %) shown in Table 2.

[0096] Furthermore, plating and / or alloying treatments were appropriately applied while the cold-rolled steel sheets were cooled to the cooling stop temperature shown in Table 2. In the "Plating Type" column in Table 2, "None" means no plating treatment was applied, GI means hot-dip galvanizing was applied, GA means alloyed hot-dip galvanizing was applied, and EG means electro-galvanizing was applied. Finally, the final cold-rolled steel sheets or plated steel sheets were obtained by reheating and holding under the conditions shown in Table 2 to create a metallic structure.

[0097] [Table 1]

[0098] [Table 2-1]

[0099] [Table 2-2]

[0100] The properties of the obtained steel plates were measured and evaluated by the following method.

[0101] [Tensile strength (TS) and uniform elongation (uEL)] The tensile strength (TS) and uniform elongation (uEL) were determined by taking a JIS No. 5 test specimen from the orientation where the longitudinal direction of the test specimen is parallel to the direction perpendicular to the rolling direction of the steel plate (direction C), and performing a tensile test in accordance with JIS Z 2241:2022.

[0102] [Work hardening ability in the later stages of deformation] The work hardening ability in the later stages of deformation was evaluated as follows. Specifically, the region in which the nominal strain during tensile deformation is 0.1 or higher, when the same tensile test as in the TS and uEL measurements is performed, was simulated as the later stages of deformation such as press forming. The rate of decrease of the instantaneous n value (rate of decrease of instantaneous n value = instantaneous n value / dε) between the nominal strain of 0.1 and the strain corresponding to uEL within that region was calculated. The higher the calculated rate of decrease of the instantaneous n value, the more suppressed the decrease in work hardening ability is considered to be, and therefore it can be judged that the work hardening ability in the later stages of deformation is improved.

[0103] In this embodiment, steel sheets with a TS of 980 MPa or higher, a TS × uEL of 14140 MPa·% or higher, and a decrease rate of the instantaneous n value of -1.00 or higher were evaluated as having high strength while also exhibiting improved ductility and work hardening ability in the later stages of deformation. The results are shown in Table 3. In Table 3, "soft α" means soft ferrite, "residual γ" means retained austenite, "fM" means fresh martensite, and "tM,B" means the sum of tempered martensite and bainite.

[0104] [Table 3-1]

[0105] [Table 3-2]

[0106] Referring to Tables 1-3, it is believed that in Comparative Examples 2 and 24, because the primary heat treatment process was not performed, the microstructure in the steel sheets before the secondary heat treatment process could not be composed mainly of martensite. As a result, retained austenite with the desired X-ray diffraction peak characteristics could not be obtained in the final microstructure, and ductility and work hardening ability in the later stages of deformation were reduced.

[0107] In Comparative Example 5, the slab heating temperature during the heating process was less than 1230°C, which prevented sufficient diffusion of Mn within the slab, resulting in a non-uniform Mn concentration distribution within the slab. Consequently, even after subsequent heat treatment, the amount of Mn distributed within the retained austenite remained non-uniform, making it impossible to obtain retained austenite with the desired X-ray diffraction peak characteristics. This resulted in reduced ductility and work hardening ability in the later stages of deformation.

[0108] In Comparative Example 6, the maximum heating temperature in the primary heat treatment process was below the Ac3 point, resulting in insufficient austenitization. It is believed that subsequent cooling failed to establish a martensite-based microstructure in the steel sheet. Consequently, the final resulting microstructure failed to obtain retained austenite with the desired X-ray diffraction peak characteristics, leading to reduced ductility and work hardening ability in the later stages of deformation.

[0109] In Comparative Example 7, the maximum heating temperature in the secondary heat treatment step was higher than (Ac3-10)°C, which led to excessive austenitization. As a result, the lath-like structure of martensite obtained in the primary heat treatment step was not retained, and the reverse transformation from the prior austenite grains could not proceed sufficiently, preventing the formation of sufficient needle-shaped retained austenite. Consequently, retained austenite with the desired X-ray diffraction peak characteristics could not be formed, and ductility and work hardening ability in the later stages of deformation were reduced.

[0110] In Comparative Example 10, the holding time during the heating process was less than 1.5 hours, which prevented sufficient diffusion of Mn within the slab, resulting in a non-uniform Mn concentration distribution within the slab. Consequently, even after subsequent heat treatment, the amount of Mn distributed within the retained austenite remained non-uniform, making it impossible to obtain retained austenite with the desired X-ray diffraction peak characteristics. This resulted in reduced ductility and work hardening ability in the later stages of deformation.

[0111] In Comparative Example 11, the value of the left-hand side of Equation 2 during the heating process was greater than 0.70, which is thought to have resulted in a non-uniform Mn concentration distribution in the slab. As a result, the amount of Mn distributed in the retained austenite remained non-uniform even after subsequent heat treatment, making it impossible to obtain retained austenite with the desired X-ray diffraction peak characteristics, and resulting in reduced ductility and work hardening ability in the later stages of deformation.

[0112] In Comparative Example 12, the cooling stop temperature in the primary heat treatment process exceeded 300°C, which is thought to have caused ferrite formation during cooling, preventing the martensite area ratio in the steel sheet from reaching 90% or more. As a result, the final obtained metal structure had a high ferrite area ratio, leading to a decrease in TS. Furthermore, it was not possible to obtain sufficient retained austenite with the desired X-ray diffraction peak characteristics, resulting in reduced ductility and work hardening ability in the later stages of deformation.

[0113] In Comparative Examples 14 and 15, the holding time in the primary heat treatment process was less than 10 seconds, resulting in insufficient austenitization. It is believed that subsequent cooling failed to establish a martensite-based microstructure in the steel sheets. Consequently, the final microstructure failed to yield retained austenite with the desired X-ray diffraction peak characteristics, leading to reduced ductility and work hardening ability in the later stages of deformation.

[0114] In Comparative Example 20, the reheating temperature in the secondary heat treatment process exceeded 450°C, resulting in excessive tempering of the martensite and the generation of a relatively large amount of carbides, making it impossible to achieve the desired area ratio of retained austenite. As a result, ductility and work hardening ability in the later stages of deformation were reduced.

[0115] In Comparative Example 21, the average cooling rate in the 500-700°C temperature range during the primary heat treatment process was less than 10°C / second. This likely resulted in the formation of ferrite during cooling, preventing the martensite area ratio in the steel sheet from reaching 90% or more. Consequently, the final resulting microstructure had a high ferrite area ratio, leading to a decrease in TS (Total Score).

[0116] In Comparative Example 30, the reheating temperature in the secondary heat treatment process was less than 300°C, which prevented the tempering of the martensite from being promoted, and thus prevented carbon enrichment into the untransformed austenite. As a result, the desired area ratio of retained austenite could not be achieved in the final resulting microstructure, and ductility and work hardening ability in the later stages of deformation were reduced.

[0117] In Comparative Example 36, the holding time at the maximum heating temperature during the secondary heat treatment process was less than 5 seconds, which is thought to have resulted in insufficient cementite dissolution in the steel and insufficient austenite formation. As a result, the area ratio of retained austenite obtained after the secondary heat treatment decreased, leading to a decrease in ductility.

[0118] In Comparative Example 43, the holding time during the reheating process in the secondary heat treatment step was less than 100 seconds, which prevented sufficient bainite growth. This resulted in insufficient carbon enrichment into the untransformed austenite, and consequently, insufficient stabilization of the retained austenite. As a result, the desired area ratio of retained austenite could not be achieved in the final microstructure, leading to reduced ductility and work hardening ability in the later stages of deformation.

[0119] In Comparative Examples 45 and 46, it is believed that the cooling stop temperature in the secondary heat treatment process was 50°C or higher, which prevented sufficient martensitic transformation of the relatively unstable austenite with low alloy concentrations of C and Mn. As a result, the final obtained metal structure failed to form fine, needle-shaped retained austenite with relatively high concentrations of alloying elements, i.e., retained austenite with the desired X-ray diffraction peak characteristics, leading to reduced ductility and work hardening ability in the later stages of deformation.

[0120] Comparative Example 51 had a low Mn content, which led to a decrease in hardenability and the formation of a large amount of ferrite, resulting in a decrease in TS. Comparative Example 49 had a low Si content, which led to a decrease in uEL and a decrease in ductility. Comparative Example 50 had a low Al content, which resulted in a low area ratio of retained austenite, leading to a decrease in ductility and work hardening ability in the later stages of deformation. Comparative Example 52 had a high Mn content, which resulted in the formation of a relatively large amount of hard phase, resulting in a decrease in ductility. In Comparative Example 53, the value of the middle side of Equation 1 was less than 1.00, which is thought to have caused excessive segregation of Si and Al at the prior austenite grain boundaries in the martensite structure before secondary heat treatment, and a significant decrease in the amount of austenite that underwent reverse transformation from the prior austenite grain boundaries during secondary heat treatment. Related to this, the area ratio of hard phases such as fresh martensite, tempered martensite, and bainite obtained after secondary heat treatment decreased, the area ratio of ferrite increased, and the TS decreased.

[0121] In contrast, in all the examples (inventive examples) of steel sheets, by appropriately controlling each condition in the manufacturing method, the microstructure of a steel sheet having a predetermined chemical composition was controlled to have a Vickers hardness of 300 Hv or more, and the microstructure was composed of a predetermined amount of retained austenite, soft ferrite, and fresh martensite. Furthermore, the morphology of retained austenite observed by X-ray diffraction was appropriately controlled. More specifically, the ratio of the integrated intensity / peak height of the X-ray diffraction peak at the (220) plane of the fcc phase at a position 1 / 4 of the thickness from the steel sheet surface was controlled to 0.100 deg. or less, and the ratio of the integrated intensity / peak height of the X-ray diffraction peak at the (311) plane of the fcc phase at a position 1 / 4 of the thickness from the steel sheet surface was controlled to 0.120 deg. or less. As a result, excellent ductility and work hardening ability in the later stages of deformation were achieved despite the TS being 980 MPa or more.

Claims

1. The chemical composition is expressed in mass percent. C: 0.050% to 0.500%, Si: 0.20% to 2.00%, Mn: 1.00% to 4.00%, P: 0.1000% or less, S: 0.0100% or less, Al: 0.200% to 1.500%, N: 0.0150% or less, O: 0.0100% or less, Ti: 0-0.300%, V: 0-1.00%, Nb: 0 to 0.100%, Cr: 0-2.00%, Ni: 0-2.000%, Cu: 0-2.000%, Co: 0-2.00%, Mo: 0-1.00%, W: 0-1.00%, B: 0 to 0.0100%, Ta: 0 to 1.00%, Sn: 0-1.00%, Sb: 0 to 1.00%, Ca: 0-0.0500%, Mg: 0 to 0.0500%, Zr: 0 to 0.5000%, Hf: 0-0.0100%, Bi: 0 to 0.0100%, REM: 0-0.0100%, As: 0-0.100%, and The remainder consists of Fe and impurities, and satisfies the following formula 1. 1.00≦Mn / (Al+0.5Si)≦4.00...Formula 1 The metallic structure, in area percentage, Soft ferrite: 20-70%, Residual austenite: 5-35% Fresh martensite: 8-35%, and Total amount of tempered martensite and bainite: 0-60%, The ratio of the integrated intensity / peak height of the X-ray diffraction peak in the (220) plane of the fcc phase at a position 1 / 4 of the plate thickness from the surface is 0.100 deg. or less, and The ratio of the integrated intensity / peak height of the X-ray diffraction peak in the (311) plane of the fcc phase at a position 1 / 4 of the plate thickness from the surface is 0.120 deg. or less. A steel plate characterized by having a Vickers hardness of 300 Hv or higher.

2. The aforementioned chemical composition is, in mass%, N: 0.0100% or less, Ca: 0-0.0100%, Mg: 0-0.0100%, and Zr: 0~0.0100% The steel plate according to claim 1, characterized by including the following:

3. The aforementioned chemical composition is, in mass%, Ti: 0.001 to 0.300%, V: 0.001-1.00%, Nb: 0.001 to 0.100%, Cr: 0.001-2.00%, Ni: 0.001 to 2.000%, Cu: 0.001-2.000%, Co: 0.001 to 2.00%, Mo: 0.001-1.00%, W: 0.001-1.00%, B: 0.0001 to 0.0100%, Ta: 0.001 to 1.00%, Sn: 0.001 to 1.00%, Sb: 0.001 to 1.00%, Ca: 0.0001-0.0500%, Mg: 0.0001-0.0500%, Zr: 0.0001 to 0.5000%, Hf: 0.0001-0.0100%, Bi: 0.0001 to 0.0100%, and REM: 0.0001-0.0100%, As: 0.001-0.100% The steel plate according to claim 1 or 2, characterized in that it includes at least one of the following.

4. The aforementioned metallographic structure, in area %, A steel sheet according to claim 1 or 2, characterized by containing 12 to 35% fresh martensite.

5. A component characterized by comprising the steel plate described in claim 1 or 2.

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