High-strength steel plates and their manufacturing methods

Abstract

The object of this invention is to provide a high-strength steel sheet with a strength (TS) of 980 MPa or higher, excellent ductility, porosity, and bendability, and whose ductility does not decrease after plating treatment, and a method for manufacturing the same. A high-strength steel sheet has a specified composition, comprising a steel microstructure with ferrite of 1% to 40% by area ratio, fresh martensite of 1% to 20% by area ratio, bainite and tempered martensite totaling 35% to 90% by area ratio, and retained austenite of 6% or higher. The average Mn content (mass%) in the retained austenite divided by the average Mn content (mass%) in the ferrite yields a value of 1.1 or higher, and the average C content (mass%) in the retained austenite with an aspect ratio of 2.0 or higher divided by the average C content (mass%) in the ferrite yields a value of 3.0 or higher. The total C content in the retained austenite divided by the C content in the T0 composition yields a value of 1.0 or higher.

Description

Technical Field

[0001] This invention relates to high-strength steel sheets with excellent formability suitable for use as components in industries such as automobiles and electrical equipment, and to a method for manufacturing such sheets. In particular, it is desirable to obtain high-strength steel sheets with a tensile strength of TS (tensile strength) of 980 MPa or higher, excellent ductility, and excellent hole expansion and bending properties. Background Technology

[0002] In recent years, improving the fuel efficiency of automobiles has become an important issue from the perspective of protecting the Earth's environment. Therefore, there is an increasing trend towards achieving thinner walls and lighter vehicle bodies through high-strength body materials. On the other hand, increasing the strength of steel sheets leads to a decrease in formability, thus there is a desire to develop materials that combine high strength and high formability.

[0003] As a high-strength steel sheet with excellent high ductility, a high-strength steel sheet with processing-induced phase transformation utilizing retained austenite has been proposed. Such a steel sheet exhibits a microstructure with retained austenite, which is easily utilized during the forming process. Furthermore, the retained austenite undergoes martensitization after forming, thus possessing high strength.

[0004] For example, Patent Document 1 proposes a high-strength steel sheet with a tensile strength of 1000 MPa or more and a total elongation (EL) of 30% or more, utilizing processing-induced phase transformation of retained austenite, exhibiting extremely high ductility. Such a steel sheet is manufactured by austenitizing a steel sheet with C, Si, and Mn as its basic components, followed by isothermal quenching within the bainitic transformation temperature range and isothermal holding—a process known as isothermal quenching. This isothermal quenching process enriches C into austenite, generating retained austenite. However, to obtain a large amount of retained austenite, a significant amount of C exceeding 0.3% needs to be added. However, as the C concentration in the steel increases, spot weldability decreases, especially at C concentrations exceeding 0.3%, where the decrease is significant, making it difficult to practically use as automotive steel. Furthermore, the aforementioned patent document prioritizes improving the ductility of high-strength thin steel sheets, and therefore does not consider hole expansion properties.

[0005] Patent Document 2 uses steel containing 4% to 6% Mn by weight and performs heat treatment in the dual-phase region of ferrite and austenite, thereby achieving a high strength-ductility balance. However, Patent Document 2 does not investigate the improvement in ductility resulting from the enrichment of Mn into the untransformed austenite, leaving room for improvement in workability.

[0006] Furthermore, Patent Document 3 discloses a method of heat treatment in the ferrite-austenite dual-phase region using steel containing 3.0% to 7.0% by mass of Mn. As a result, the total elongation is increased by enriching Mn into the untransformed austenite, forming stable retained austenite. However, due to the short heat treatment time and slow Mn diffusion rate, it is speculated that the enrichment of Mn is insufficient to achieve not only elongation but also porosity and flexibility.

[0007] Furthermore, Patent Document 4 discloses a method of using steel containing 0.50% to 12.00% by mass of Mn to perform long-term heat treatment on hot-rolled plates in the ferrite-austenite dual-phase region. As a result, retained austenite with a large aspect ratio is formed, promoting the enrichment of Mn into the untransformed austenite, thus improving uniform elongation. However, no research was conducted on improving hole expansion properties, bending performance, and elongation simultaneously. During plating and alloying processes, austenite is prone to decomposition, making it difficult to ensure the required amount of retained austenite.

[0008] Existing technical documents

[0009] Patent documents

[0010] Patent Document 1: Japanese Patent Application Publication No. 61-157625

[0011] Patent Document 2: Japanese Patent Application Publication No. 1-259120

[0012] Patent Document 3: Japanese Patent Application Publication No. 2003-138345

[0013] Patent Document 4: Japanese Patent No. 6123966 Summary of the Invention

[0014] The problem that the invention aims to solve

[0015] This invention was made in view of the current situation as described above, and its object is to provide a high-strength steel sheet with a tensile strength (TS) of 980 MPa or higher, excellent formability, and no decrease in ductility after plating treatment, as well as a method for manufacturing the same. Formability as described herein refers to ductility, hole-expanding ability, and bending ability.

[0016] Methods for solving problems

[0017] In order to solve the above problems and to manufacture high-strength steel plates with excellent formability, the inventors have conducted repeated and in-depth research from the perspective of the composition of the steel plates and the manufacturing method, and have discovered the following:

[0018] It was found that a composition containing 2.00% to 8.00% by mass of Mn, with appropriate adjustments to other alloying elements such as Ti, was hot-rolled, held at a temperature below the Ac1 transformation point for more than 1800 seconds as needed, and then pickled as required before cold rolling. Next, it was held at a temperature above -50°C from the Ac3 transformation point for more than 20 seconds but less than 1800 seconds, cooled to a cooling stop temperature below the martensitic transformation start temperature, and then reheated to a reheating temperature between 120°C and 450°C. Then, it was held at the above reheating temperature for more than 2 seconds but less than 1800 seconds, and then cooled to room temperature. This process is crucial for the formation of a film-like austenite enriched with C, which serves as the core of fine residual austenite with a large aspect ratio and significant enrichment of Mn and C, during the subsequent annealing process.

[0019] Furthermore, after the above cooling, the temperature is maintained for 20 seconds to 600 seconds within a temperature range above -20°C from the Ac1 phase transformation point, then cooled to a cooling stop temperature below the martensitic phase transformation start temperature, and then reheated to a reheating temperature range above 120°C to 480°C. Then, the temperature is maintained at the above reheating temperature for 2 seconds to 600 seconds, and then cooled to room temperature. The results show that a high-strength steel sheet with excellent formability can be manufactured. The steel sheet is characterized by having a steel microstructure with an area ratio of 1% to 40% ferrite, 1% to 20% fresh martensite, a sum of 35% to 90% bainite and tempered martensite, and 6% or more retained austenite. The average Mn content (mass%) in the retained austenite divided by the average Mn content (mass%) in the ferrite is 1.1 or more, and the average C content (mass%) in the retained austenite with an aspect ratio of 2.0 or more divided by the average C content (mass%) in the ferrite is 3.0 or more. The total C content in the retained austenite divided by the C content in the T0 composition is 1.0 or more.

[0020] This invention is based on the above insights, and its main points are as follows.

[0021] [1] A high-strength steel plate having the following characteristics:

[0022] The composition, by mass percent, comprises: C: 0.030% or more and 0.250% or less; Si: 0.01% or more and 3.00% or less; Mn: 2.00% or more and 8.00% or less; P: 0.100% or less; S: 0.0200% or less; N: 0.0100% or less; Al: 0.001% or more and 2.000% or less, with the balance being Fe and unavoidable impurities; and by area percent, ferrite comprises 1% or more and 40% or less, and fresh martensite comprises 1% or more and 20% or less. For steel structures containing 35% or less bainite and 90% or less bainite and tempered martensite, and 6% or more retained austenite, the average Mn content (mass%) in the retained austenite divided by the average Mn content (mass%) in the ferrite is 1.1 or more, and the average C content (mass%) in the retained austenite with an aspect ratio of 2.0 or more is 3.0 or more, and the C content in all retained austenite divided by the C content in the T0 composition is 1.0 or more.

[0023] [2] The high-strength steel plate according to [1], wherein the above composition, in mass %, further contains at least one element selected from Ti: less than 0.200%, Nb: less than 0.200%, V: less than 0.500%, W: less than 0.500%, B: less than 0.0050%, Ni: less than 1.000%, Cr: less than 1.000%, Mo: less than 1.000%, Cu: less than 1.000%, Sn: less than 0.200%, Sb: less than 0.200%, Ta: less than 0.100%, Zr: less than 0.200%, Ca: less than 0.0050%, Mg: less than 0.0050%, REM: less than 0.0050%.

[0024] [3] The high-strength steel plate according to [1] or [2], wherein the area ratio of blocky retained austenite divided by the area ratio of all retained austenite and blocky fresh martensite is less than 0.5.

[0025] [4] The high-strength steel plate described in [1] to [3] further has a galvanized layer on its surface.

[0026] [5] The high-strength steel plate according to [4], wherein the zinc coating is an alloyed zinc coating.

[0027] [6] A method for manufacturing a high-strength steel plate, which is the method for manufacturing a high-strength steel plate as described in any one of [1] to [3], wherein a steel billet having the composition described in [1] or [2] is heated, hot-rolled at a temperature of 750°C or higher and 1000°C or lower at the exit temperature of the finishing mill, coiled and cold-rolled at a temperature of 300°C or higher and 750°C or lower, and then cooled to a temperature range of -50°C or higher from the Ac3 phase transformation point for 20s or more and 1800s or less, and then cooled to a temperature below the martensitic phase transformation start temperature. After reheating to a temperature between 120°C and 450°C, the temperature is held at the reheating temperature for 2 seconds to 1800 seconds, then cooled to room temperature. After holding at a temperature between -20°C and 600°C above the Ac1 phase transformation point, the temperature is cooled to a cooling stop temperature below the martensitic phase transformation start temperature. After reheating to a temperature between 120°C and 480°C, the temperature is held at the reheating temperature for 2 seconds to 600 seconds, then cooled to room temperature.

[0028] [7] The method for manufacturing high-strength steel plates according to [6] further includes galvanizing.

[0029] [8] According to the manufacturing method of high-strength steel plate described in [7], alloying treatment is carried out at a temperature above 450°C and below 600°C after the above-mentioned galvanizing treatment.

[0030] [9] The method for manufacturing high-strength steel sheet according to any one of [6] to [8], wherein, after the above-mentioned coiling and before cold rolling, the temperature range below the Ac1 phase transformation point is maintained for more than 1800 s.

[0031] Invention Effects

[0032] According to the present invention, a high-strength steel sheet with a tensile strength (TS) of over 980 MPa, excellent formability after plating, and particularly excellent ductility, hole expansion, and bending properties, and whose ductility is not reduced after plating, can be obtained. By applying the high-strength steel sheet obtained using the manufacturing method of the present invention to, for example, automotive structural components, the improved fuel efficiency resulting from vehicle body lightweighting can be achieved, demonstrating significant industrial value. Detailed Implementation

[0033] The present invention will now be described in detail. It should be noted that the "%" indicating the content of the constituent elements refers to "mass %" unless otherwise stated.

[0034] (1) The reasons for limiting the composition of steel to the above range in this invention will be explained.

[0035] C: Above 0.030% and below 0.250%

[0036] Carbon (C) is an element required to increase strength by enabling the formation of martensite and other phase transformations at low temperatures. Furthermore, C is effective in improving the stability of retained austenite and enhancing the ductility of steel. When the C content is less than 0.030%, excessive ferrite is formed, thus failing to achieve the desired strength. Additionally, it is difficult to ensure a sufficient area ratio of retained austenite, resulting in poor ductility. On the other hand, when C exceeds 0.250% and is excessively present, the area ratio of hard martensite becomes too large. During porosity testing, micro-voids at the martensite grain boundaries increase, and crack propagation occurs, reducing porosity. Furthermore, significant hardening occurs in the weld and heat-affected zone, reducing the mechanical properties of the weld and thus deteriorating spot weldability and arc weldability. From this perspective, the C content is set to 0.030% or more and 0.250% or less. A lower limit of 0.080% or more is preferred. An upper limit of 0.200% or less is also preferred.

[0037] Si: ≥0.01% and ≤3.00%

[0038] Si enhances the work hardening ability of ferrite, thus effectively ensuring good ductility. When the Si content is less than 0.01%, its effect is insufficient; therefore, a lower limit is set at 0.01%. However, excessive addition of Si exceeding 3.00% leads to decreased ductility and flexibility due to steel embrittlement, and further deteriorates surface properties due to the formation of red oxide scale, etc., further reducing coating quality. Therefore, Si is set between 0.01% and 3.00%. A lower limit of 0.20% or more is preferred. An upper limit of 2.00% or less is preferred, and more preferably less than 1.20%.

[0039] Mn: ≥2.00% and ≤8.00%

[0040] Mn is an extremely important element in this invention. Mn stabilizes retained austenite, is effective in ensuring good ductility, and increases the strength of steel through solid solution strengthening. This effect is confirmed when the Mn content in the steel is 2.00% or more. However, excessive addition of Mn exceeding 8.00% degrades chemical conversion treatability and coating quality. From this perspective, the Mn content is set to 2.00% or more and 8.00% or less. A preferred lower limit is 2.30% or more, more preferably 2.50% or more. Furthermore, a preferred upper limit is 6.00% or less, more preferably 4.20% or less.

[0041] P: below 0.100%

[0042] P is an element that provides solid solution strengthening and can be added according to the desired strength. When the amount of P exceeds 0.100%, it leads to deterioration of weldability and reduces the alloying rate during alloying treatment of the galvanized layer, thus impairing the quality of the galvanized layer. The lower limit can be 0%, but considering production costs, it is preferable to be 0.001% or more. Therefore, the amount of P is set to 0.100% or less. A more preferred lower limit is 0.005% or more. A preferred upper limit is set to 0.050% or less.

[0043] S: below 0.0200%

[0044] S-axis grain boundary segregation causes steel to become embrittled during hot working and, existing in the form of sulfides, reduces local deformation capacity. Therefore, its amount needs to be set to 0.0200% or less, preferably 0.0100% or less, and more preferably 0.0050% or less. The lower limit can be 0%, but from the perspective of production costs, it is preferable to be 0.0001% or more.

[0045] N: below 0.0100%

[0046] Nitrogen (N) is an element that degrades the aging resistance of steel. In particular, the degradation of aging resistance becomes significant when the N content exceeds 0.0100%. Lower amounts are preferred; the lower limit can be 0%, but considering production costs, the N content is preferably 0.0005% or more. Therefore, the N content is set to 0.0100% or less. More preferably, it is set to 0.0010% or more. The upper limit of the N content is preferably set to 0.0070% or less.

[0047] Al: Above 0.001% and below 2.000%

[0048] Al is an element that expands the two-phase region between ferrite and austenite, reduces the dependence of mechanical properties on annealing temperature, and is effective in improving material stability. When the Al content is less than 0.001%, its effect is insufficient; therefore, a lower limit is set at 0.001%. Furthermore, Al acts as a deoxidizer and is effective in improving the cleanliness of steel; it is preferably added during the deoxidation process. However, excessive addition exceeding 2.000% increases the risk of steel sheet breakage during continuous casting, reducing manufacturability. From this perspective, the Al content is set to 0.001% or more and 2.000% or less. A preferred lower limit is 0.025% or more, more preferably 0.200% or more. A preferred upper limit is 1.200% or less.

[0049] In addition, based on the above composition, it may contain, by mass%, at least one element selected from Ti: 0.200% or less, Nb: 0.200% or less, V: 0.500% or less, W: 0.500% or less, B: 0.0050% or less, Ni: 1.000% or less, Cr: 1.000% or less, Mo: 1.000% or less, Cu: 1.000% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ta: 0.1000% or less, Zr: 0.200% or less, Ca: 0.0050% or less, Mg: 0.0050% or less, REM: 0.0050% or less.

[0050] Ti: below 0.200%

[0051] Ti is effective in precipitation strengthening of steel. By increasing the strength of ferrite, it can reduce the hardness difference with the hard second phase (martensite or retained austenite), ensuring better porosity. Therefore, it can be included as needed. However, when it exceeds 0.200%, the area fraction of hard martensite becomes too large. During porosity expansion tests, the micropores at the martensite grain boundaries increase, and crack propagation occurs, sometimes reducing porosity. Therefore, when adding Ti, its addition amount is set to 0.200% or less. The preferred lower limit is 0.005% or more, more preferably 0.010% or more. The preferred upper limit is set to 0.100% or less.

[0052] Nb: below 0.200%, V: below 0.500%, W: below 0.500%

[0053] Nb, V, and W are effective in precipitation strengthening of steel, similar to the effect of Ti addition. By increasing the strength of ferrite, they can reduce the hardness difference with the hard second phase (martensite or retained austenite), ensuring better porosity. Therefore, they can be included as needed. However, when Nb exceeds 0.200% and V and W exceed 0.500%, the area fraction of hard martensite becomes too large. During porosity expansion tests, micro-voids at the martensite grain boundaries increase, and crack propagation occurs, sometimes reducing porosity. Therefore, when Nb is added, its addition amount is set to 0.200% or less. Preferably, the lower limit of Nb is set to 0.005% or more, more preferably 0.010% or more. Preferably, the upper limit of Nb is set to 0.100% or less. When V and W are added, their addition amounts are each set to 0.500% or less. Preferably, the lower limit of V and W is set to 0.005% or more, more preferably 0.010% or more. The upper limits of V and W are preferably set to be below 0.300%.

[0054] B: Below 0.0050%

[0055] Boron (B) inhibits the formation and growth of ferrite from austenite grain boundaries. By increasing the strength of ferrite, it reduces the hardness difference with the hard second phase (martensite or retained austenite), ensuring better porosity. Therefore, it can be included as needed. However, when it exceeds 0.0050%, formability sometimes decreases. Therefore, when adding B, its addition amount is set to 0.0050% or less. The preferred lower limit is set to 0.0003% or more, more preferably 0.0005% or more. The preferred upper limit is set to 0.0030% or less.

[0056] Ni: below 1.000%

[0057] Ni is an element that stabilizes retained austenite, effectively ensuring better ductility, and further increases the strength of steel through solid solution strengthening. Therefore, it can be included as needed. On the other hand, when added at more than 1.000%, the area ratio of hard martensite becomes excessive. During porosity expansion tests, the micro-voids at the martensite grain boundaries increase, and crack propagation occurs, reducing porosity. Therefore, when adding Ni, its addition amount is set to 1.000% or less, preferably 0.005% or more and 1.000% or less.

[0058] Cr: less than 1.000%, Mo: less than 1.000%

[0059] Cr and Mo have the effect of improving the balance between strength and ductility, and therefore can be added as needed. However, when Cr and Mo are added in excess, exceeding 1.000% for Cr and 1.000% for Mo respectively, the area ratio of hard martensite becomes too large. During porosity expansion tests, the micropores at the grain boundaries of martensite increase, and crack propagation occurs, sometimes reducing porosity. Therefore, when adding these elements, their amounts are set to 1.000% or less for Cr and 1.000% or less for Mo, preferably 0.005% or more but less than 1.000% for Cr and 0.005% or more but less than 1.000%.

[0060] Cu: below 1.000%

[0061] Cu is an effective element for strengthening steel, and can be used to strengthen steel as needed, as long as it is within the range specified in this invention. On the other hand, when added in amounts exceeding 1.000%, the area ratio of hard martensite becomes excessive, leading to an increase in micropores at the grain boundaries of martensite during porosity expansion tests, and the propagation of cracks, resulting in reduced porosity. Therefore, when adding Cu, its amount is set to 1.000% or less, preferably 0.005% or more and 1.000% or less.

[0062] Sn: less than 0.200%, Sb: less than 0.200%

[0063] From the viewpoint of suppressing decarburization in a region of approximately tens of μm on the surface of the steel sheet caused by nitriding and oxidation, Sn and Sb are added as needed. Suppressing such nitriding and oxidation, and preventing a reduction in the martensite area ratio on the steel sheet surface, is effective in ensuring strength and material stability; therefore, they can be included as needed. On the other hand, excessive addition of any of these elements exceeding 0.200% will lead to a decrease in toughness. Therefore, when adding Sn and Sb, their contents are each set to 0.200% or less, preferably 0.002% or more and 0.200% or less.

[0064] Ta: below 0.100%

[0065] Like Ti and Nb, Ta contributes to high strength by forming alloy carbides and alloy carbonitrides. Furthermore, Ta is believed to have the following effect: it partially dissolves in Nb carbides and Nb carbonitrides to form composite precipitates such as (Nb, Ta) and (C, N), thereby significantly suppressing precipitate coarsening and stabilizing the contribution of precipitation strengthening to strength. Therefore, Ta can be included as needed. On the other hand, even with excessive Ta addition, the precipitate stabilization effect becomes saturated, and the alloy cost increases. Therefore, when Ta is added, its content is set to 0.100% or less, preferably 0.001% or more and 0.100% or less.

[0066] Zr: below 0.200%

[0067] Zr is an effective element for shaping sulfides into spherical forms and improving their adverse effects on flexibility; therefore, it can be included as needed. However, excessive addition exceeding 0.200% can lead to an increase in inclusions and other defects, causing surface and internal defects. Therefore, when adding Zr, the amount added is set to 0.200% or less, preferably 0.0005% or more and 0.200% or less.

[0068] Ca: less than 0.0050%, Mg: less than 0.0050%, REM: less than 0.0050%

[0069] Ca, Mg, and REM are effective elements for shaping sulfides into spherical forms and mitigating their adverse effects on porosity; therefore, they can be included as needed. However, excessive addition of each exceeding 0.0050% can lead to an increase in inclusions and other defects, causing surface and internal defects. Therefore, when adding Ca, Mg, and REM, the amount added is set to 0.0050% or less, preferably 0.0005% or more and 0.0050% or less.

[0070] The balance other than the above components is Fe and unavoidable impurities.

[0071] (2) Next, the steel structure will be described.

[0072] Ferrite area fraction: 1% or more and 40% or less

[0073] To ensure sufficient ductility, the ferrite area fraction needs to be 1% or more. Furthermore, to ensure a strength tolerance (TS) of 980 MPa or higher, the soft ferrite area fraction needs to be 40% or less. It should be noted that the ferrite referred to here is polygonal ferrite, granular ferrite, or acicular ferrite, which is a relatively soft and ductile ferrite. Preferably, it is 3% or more and 30% or less.

[0074] Area percentage of fresh martensite: 1% or more and 20%

[0075] To achieve a strength TS of 980 MPa or higher, the area fraction of fresh martensite needs to be 1% or more. Furthermore, to ensure good porosity, the area fraction of fresh martensite needs to be 20% or less. Preferably, it is 3% or more and 18% or less.

[0076] The sum of the area ratios of bainite and tempered martensite is 35% or more and 90% or less.

[0077] Bainite and tempered martensite are effective microstructures for improving porosity. When the sum of the area ratios of bainite and tempered martensite is less than 35%, good porosity is not achieved. Therefore, the sum of the area ratios of bainite and tempered martensite needs to be 35% or more. On the other hand, when the sum of the area ratios of bainite and tempered martensite is greater than 90%, retained austenite, which is expected to provide ductility, is not obtained, and therefore good ductility is not achieved. Therefore, the sum of the area ratios of bainite and tempered martensite needs to be 90% or less. Preferably, it is 45% or more and 85% or less.

[0078] It should be noted that the area ratios of ferrite, fresh martensite, tempered martensite, and bainite can be calculated as follows: After grinding a section (L section) of the plate thickness parallel to the rolling direction of the steel plate, it is etched with a 3% (v / v) nitric acid ethanol solution. At the 1 / 4 thickness position (equivalent to 1 / 4 of the plate thickness from the surface in the depth direction), 10 fields of view are observed using a SEM (scanning electron microscope) at 2000x magnification. Using the obtained microstructure images, the area ratios of each microstructure (ferrite, fresh martensite, tempered martensite, and bainite) in the 10 fields of view are calculated using MediaCybernetics' Image-Pro, and these values ​​are averaged. Furthermore, in the above microstructure images, ferrite appears as a gray microstructure (base microstructure), fresh martensite appears as a white microstructure, tempered martensite appears as a microstructure with a gray internal structure within white martensite, and bainite appears as a dark gray microstructure with numerous linear grain boundaries.

[0079] Area ratio of retained austenite: 6% or more

[0080] To ensure sufficient ductility, the area fraction of retained austenite needs to be 6% or more. Preferably, it is 8% or more. More preferably, it is 10% or more.

[0081] It should be noted that, regarding the area ratio of retained austenite, after grinding the steel plate to a surface 0.1 mm away from 1 / 4 of the plate thickness, another 0.1 mm was removed by chemical grinding. For the surface obtained in this way, the integral intensity ratio of the diffraction peaks of the {200}, {220}, and {311} planes of fcc iron and the {200}, {211}, and {220} planes of bcc iron was determined by using an X-ray diffraction device with CoKα rays. The nine integral intensity ratios obtained were then averaged to obtain the area ratio.

[0082] The value obtained by dividing the average Mn content (mass%) in the retained austenite by the average Mn content (mass%) in the ferrite is 1.1 or higher.

[0083] A value of 1.1 or higher, obtained by dividing the average Mn content (mass%) in the retained austenite by the average Mn content (mass%) in the ferrite, is an extremely important constituent element in this invention. To ensure good ductility, a high area ratio of stable retained austenite enriched with Mn is required. Preferably, it is 1.2 or higher.

[0084] The value obtained by dividing the average carbon content (mass%) in the retained austenite with an aspect ratio of 2.0 or higher by the average carbon content (mass%) in the ferrite is 3.0 or higher.

[0085] A value of 3.0 or higher, obtained by dividing the average carbon content (mass%) in the retained austenite by the average carbon content (mass%) in the ferrite, with an aspect ratio (major axis / minor axis) of 2.0 or higher, is an extremely important constituent element in this invention. To ensure good bending properties, a high area ratio of stable, carbon-enriched retained austenite is required. Preferably, it is 5.0 or higher. It should be noted that there is no particular upper limit to the aspect ratio of the retained austenite, and it can preferably be 20.0 or lower.

[0086] The C and Mn content in retained austenite and ferrite can be quantified using FE-EPMA (Field Emission-Electron Probe Micro Analyzer) to measure the Mn distribution in each phase of the rolled section at 1 / 4 of the plate thickness. The Mn content is calculated by averaging the C and Mn content analysis results of 30 retained austenite grains and 30 ferrite grains.

[0087] To identify retained austenite from retained austenite and martensite, the same field of view was observed using SEM (Scanning Electron Microscope) and EBSD (Electron Backscattered Diffraction). Then, the retained austenite in the SEM image was identified using the EBSD phase map. It should be noted that the aspect ratio of the retained austenite was calculated by drawing an ellipse circumscribed to the retained austenite grain using Photoshop elements13 and dividing its major axis by its minor axis.

[0088] The value obtained by dividing the total carbon content in the retained austenite by the carbon content in the T0 composition is 1.0 or higher.

[0089] A value of 1.0 or higher, obtained by dividing the carbon content in all retained austenite by the carbon content in the T0 composition, is an extremely important structural element in this invention. The T0 composition is the composition in which the free energies of fcc and bcc are equal at any temperature, where austenite is fcc and ferrite and bainite are bcc. By ensuring that the carbon content in all retained austenite is higher than the carbon content in the T0 composition where the free energies of fcc and bcc are equal, the decomposition of retained austenite during plating can be suppressed, and the desired amount of retained austenite can be obtained. As a result, the reduction in ductility caused by plating in the past can be prevented, and good ductility can be ensured. Therefore, it is necessary that the value obtained by dividing the carbon content in all retained austenite by the carbon content in the T0 composition is 1.0 or higher. Preferably, it is set to 1.1 or higher.

[0090] The amount of C in the total residual austenite here is calculated using an X-ray diffraction apparatus, using Kα rays with Co, using the displacement of the diffraction peaks of the (220) plane and the following [1], [2] equations.

[0091] a=1.7889×√2 / sinθ…[1]

[0092] a=3.578+0.033[C]+0.00095[Mn]…[2]

[0093] In [1] and [2], a is the lattice constant of austenite (), and θ is the value (rad) obtained by dividing the diffraction peak angle of the (220) plane by 2. In [2], [M] is the mass percentage of element M in all austenite. In this invention, the mass percentage of element M in the retained austenite is set to the mass percentage in the whole steel.

[0094] Furthermore, the carbon content in the T0 composition can be uniquely calculated using the comprehensive thermodynamic calculation software Thermo-Calc and the database TCFE7, based on the steel's composition and its content. The calculated T0 composition is based on the reheat temperature before immersion in the galvanizing bath.

[0095] In addition, it is preferable that the value obtained by multiplying the average Mn content (mass%) in the retained austenite by the average Mn content (mass%) in the ferrite by the average aspect ratio of the retained austenite is 3.0 or higher. To ensure good ductility, it is necessary to have a high aspect ratio and a high area ratio of stable retained austenite enriched with Mn. Preferably, it is 4.0 or higher. Furthermore, a preferred upper limit is 20.0 or lower.

[0096] Furthermore, it is preferable that the area ratio of the bulky retained austenite divided by the area ratio of all retained austenite and the bulky fresh martensite is 0.5 or less. The bulky retained austenite has high stability due to the constraint from the surrounding grains; therefore, during punching, the martensitic transformation occurs in the high-strain region, increasing the hardness difference with the surrounding grains and sometimes deteriorating the porosity. Therefore, it is preferable that the area ratio of the bulky retained austenite divided by the area ratio of all retained austenite and the bulky fresh martensite is 0.5 or less. Preferably, it is 0.4 or less. It should be noted that bulky retained austenite refers to austenite with an aspect ratio less than 2.0. There is no limitation on the average grain size of the bulky retained austenite, but an average grain size of, for example, less than 3 μm can be considered. This average grain size can be determined by conventionally known methods, for example, by image analysis of microstructure images of the bulky retained austenite obtained using a scanning electron microscope (SEM).

[0097] In the steel microstructure of the present invention, even if carbides such as pearlite and cementite are present in the range of less than 10% by area, in addition to ferrite, fresh martensite, bainite, tempered martensite and retained austenite, the effect of the present invention will not be impaired.

[0098] The aforementioned high-strength steel sheet may further have a galvanized layer. This galvanized layer may be an alloyed galvanized layer that has undergone alloying treatment.

[0099] (3) Next, the manufacturing conditions will be explained.

[0100] Heating temperature of steel billet

[0101] While not specifically limited, the heating temperature of the steel billet is preferably set to 1100°C or higher and 1300°C or lower. Precipitates present during the billet heating stage exist as coarse precipitates in the final steel sheet and do not contribute to strength. Therefore, it is preferable to redissolve the Ti and Nb precipitates that appear during casting. Thus, the heating temperature of the steel billet is preferably set to 1100°C or higher. Furthermore, from the viewpoint of removing defects such as bubbles and segregation from the surface of the steel billet, and reducing cracks and unevenness on the steel sheet surface to achieve a smoother steel sheet surface, the heating temperature of the steel billet is preferably set to 1100°C or higher. On the other hand, when the heating temperature of the steel billet exceeds 1300°C, the loss of oxide scale increases with the increase in oxidation. Therefore, the heating temperature of the steel billet is preferably set to 1300°C or lower. More preferably, it is set to 1150°C or higher and 1250°C or lower.

[0102] To prevent macroscopic segregation, steel billets are preferably manufactured using continuous casting, but they can also be manufactured using ingot casting, thin-slab casting, etc. Furthermore, after manufacturing the steel billets, in addition to the existing method of temporarily cooling them to room temperature and then reheating them, energy-saving processes such as direct-feed rolling (where the billets are loaded into the furnace as warm plates without cooling to room temperature, or rolled immediately after slight holding at room temperature) can be used without any problems. Additionally, steel billets are normally produced into thin slabs through rough rolling; however, when the heating temperature is low, from the viewpoint of preventing defects during hot rolling, it is preferable to heat the thin slabs using a bar heater or similar method before finish rolling.

[0103] Finishing mill exit temperature of hot-rolled steel: above 750℃ and below 1000℃

[0104] Heated steel billets are hot-rolled through roughing and finishing rolling to produce hot-rolled steel plates. At this stage, when the finishing rolling temperature exceeds 1000°C, the formation of oxides (scale) increases dramatically, resulting in a rough interface between the steel matrix and the oxides, and a tendency for surface quality deterioration after pickling and cold rolling. Furthermore, the presence of residual hot-rolled scale after pickling negatively impacts ductility and porosity. Additionally, sometimes the grain size becomes excessively large, leading to a rough surface on the pressed product during processing. On the other hand, when the finishing rolling temperature is below 750°C, the rolling load increases, leading to a higher reduction rate of austenite in its non-recrystallized state, resulting in abnormally developed textures and significant in-plane anisotropy in the final product, impairing material homogeneity (material stability). Moreover, ductility itself decreases. Therefore, it is necessary to set the finishing rolling exit temperature of hot rolling to be above 750°C and below 1000°C. Preferably, it is set to above 800°C and below 950°C.

[0105] Coiling temperature after hot rolling: above 300℃ and below 750℃

[0106] When the coiling temperature after hot rolling exceeds 750°C, the ferrite grain size in the hot-rolled sheet structure becomes larger, making it difficult to ensure the desired strength of the final annealed sheet. On the other hand, when the coiling temperature after hot rolling is below 300°C, the strength of the hot-rolled sheet increases, the rolling load during cold rolling increases, or sheet shape defects occur, thus reducing productivity. Therefore, it is necessary to set the coiling temperature after hot rolling to be above 300°C and below 750°C. Preferably, it is set to above 400°C and below 650°C.

[0107] It should be noted that rough-rolled plates can also be joined together during hot rolling for continuous finishing. Alternatively, the rough-rolled plates can be temporarily coiled. Furthermore, to reduce the rolling load during hot rolling, some or all of the finishing rolling can be lubricated. Lubricated rolling is also effective from the viewpoint of achieving uniformity in steel plate shape and material properties. It should be noted that the coefficient of friction during lubricated rolling is preferably set to 0.10 or higher and 0.25 or lower.

[0108] The hot-rolled steel sheets manufactured in this way are pickled as needed. Pickling removes oxides from the surface of the steel sheet; therefore, it is preferred to perform pickling to ensure good chemical conversion properties and coating quality of the final high-strength steel sheet. Furthermore, pickling can be performed in a single operation or in multiple stages.

[0109] cold rolling

[0110] After winding, pickling is performed as needed, followed by cold rolling. There are no particular limitations on the cold rolling reduction rate, but it is preferably 5% to 60%.

[0111] It remains below the Ac1 phase transition point for more than 1800 seconds.

[0112] Holding the steel sheet for more than 1800 seconds within a temperature range below the Ac1 transformation point allows for softening after subsequent cold rolling, and this process can be performed as needed. When holding for more than 1800 seconds, Mn enriches in the austenite, and upon cooling, hard martensite and retained austenite are formed, sometimes preventing the steel sheet from softening. Conversely, holding for less than 1800 seconds fails to remove the strain from hot rolling, sometimes also preventing the steel sheet from softening.

[0113] It should be noted that the heat treatment method can be either continuous annealing or batch annealing. After the heat treatment, the material should be cooled to room temperature, but there are no specific requirements for the cooling method or rate; any of the following cooling methods can be used: furnace cooling or air cooling in batch annealing, and gas jet cooling, spray cooling, or water cooling in continuous annealing. Furthermore, conventional methods can be used when pickling is performed.

[0114] Maintain the temperature range above -50°C from the Ac3 phase transformation point for more than 20 seconds and less than 1800 seconds (corresponding to the first annealing treatment of the cold-rolled sheet in the embodiment).

[0115] When held at temperatures below the Ac3 transformation point of -50°C, Mn accumulates in the austenite and does not undergo a martensitic transformation during cooling, failing to produce a nucleus of retained austenite with a large aspect ratio. As a result, in the subsequent annealing process (corresponding to the second annealing treatment of the cold-rolled sheet in the embodiment), retained austenite forms from grain boundaries, increasing the amount of retained austenite with a small aspect ratio, failing to obtain the desired microstructure, and thus reducing porosity.

[0116] If the time limit is less than 20 seconds, sufficient recrystallization will not occur, resulting in an undesirable microstructure and reduced porosity. Furthermore, the Mn surface enrichment required to ensure the quality of subsequent coatings will not be adequately achieved.

[0117] On the other hand, when the temperature is maintained for more than 1800s, the Mn surface becomes excessively enriched, the coating quality deteriorates, and the austenite grains in the annealing process become coarser. As a result, the nuclei of the residual austenite formed during the subsequent cooling process also become coarser, and C cannot be sufficiently enriched to above the T0 composition, resulting in reduced ductility after coating.

[0118] Cooling to the cooling stop temperature below the martensitic transformation start temperature

[0119] When the cooling stop temperature exceeds the martensitic transformation start temperature, and the amount of martensite undergoing the transformation is small, all the untransformed austenite undergoes martensitic transformation during final cooling, making it impossible to obtain a nucleus of retained austenite with a large aspect ratio. As a result, in the subsequent annealing process (corresponding to the second annealing treatment of the cold-rolled sheet in the embodiment), retained austenite forms from grain boundaries, and the amount of retained austenite with a small aspect ratio increases, failing to obtain the desired microstructure, thus reducing ductility and porosity.

[0120] Preferably, the martensitic transformation start temperature is above -250℃ and below -50℃.

[0121] After reheating to a reheating temperature within the range of 120°C to 450°C, hold at the reheating temperature for 2 seconds to 1800 seconds, and then cool to room temperature.

[0122] At reheating temperatures below 120°C, carbon does not accumulate in the retained austenite formed during subsequent annealing processes, resulting in an undesirable microstructure and reduced ductility, bending properties, and post-plating ductility. At reheating temperatures above 450°C, the nuclei of retained austenite with large aspect ratios decompose, while the amount of retained austenite with small aspect ratios increases, again resulting in an undesirable microstructure and reduced ductility. Similarly, holding the heat for less than 2 seconds fails to produce nuclei of retained austenite with large aspect ratios, resulting in an undesirable microstructure and reduced ductility, bending properties, and post-plating ductility. Furthermore, holding the heat for more than 1800 seconds results in the nuclei of retained austenite with large aspect ratios decomposing, while the amount of retained austenite with small aspect ratios increases, again resulting in an undesirable microstructure and reduced ductility.

[0123] After reheating and holding for a specified time, temporarily cool to room temperature. There are no particular restrictions on the cooling method; any known method may be used.

[0124] Maintain the temperature range of -20°C above the Ac1 phase transformation point for more than 20 seconds and less than 600 seconds (corresponding to the second annealing treatment of the cold-rolled sheet in the embodiment).

[0125] Maintaining a temperature range above -20°C from the Ac1 transformation point for 20 seconds to 600 seconds is an extremely important inventive requirement in this invention. In temperatures below -20°C from the Ac1 transformation point and with a holding time of less than 20 seconds, the carbides formed during heating do not completely dissolve, making it difficult to ensure a sufficient area ratio of martensite and retained austenite, resulting in reduced strength. Preferably, the temperature is above the Ac1 transformation point. More preferably, it is above +20°C from the Ac1 transformation point and below +50°C from the Ac3 transformation point. Furthermore, when holding for more than 600 seconds, the austenite coarsens during annealing, thus the diffusion of Mn into the austenite becomes insufficient, preventing enrichment and failing to obtain a sufficient area ratio of retained austenite to ensure ductility.

[0126] Cooling to the cooling stop temperature below the martensitic transformation start temperature

[0127] When the cooling stop temperature exceeds the martensitic transformation start temperature, the amount of martensite undergoing the transformation is small, resulting in a small amount of martensite being tempered during subsequent reheating, thus failing to obtain the desired amount of tempered martensite. Preferably, the martensitic transformation start temperature is above -250°C and below -30°C.

[0128] After reheating to a reheating temperature within the range of 120°C to 480°C, hold at the above reheating temperature for 2 seconds to 600 seconds, and then cool to room temperature.

[0129] At reheating temperatures below 120°C, the fresh martensite is not tempered, resulting in an undesirable microstructure. At reheating temperatures above 480°C, the bainitic transformation is delayed, again failing to yield the desired microstructure. Furthermore, holding times of less than 2 seconds result in insufficient bainitic transformation, thus also failing to produce the desired microstructure. On the other hand, holding times exceeding 600 seconds lead to carbide precipitation during the bainitic transformation, reducing the carbon content in the retained austenite and resulting in an undesirable microstructure.

[0130] After maintaining the temperature at that level for a specified time, cool to room temperature. There are no particular limitations on the cooling method; any known method may be used.

[0131] Zinc plating

[0132] The obtained high-strength steel sheet is then galvanized as needed. In the case of hot-dip galvanizing, the steel sheet that has undergone the above-mentioned annealing treatment is immersed in a galvanizing bath at a temperature of 440°C to 500°C for hot-dip galvanizing. The coating adhesion is then adjusted by methods such as gas wiping. It should be noted that a galvanizing bath with an Al content of 0.08% to 0.30% is preferably used for hot-dip galvanizing.

[0133] When alloying the hot-dip galvanized layer, the alloying process should be carried out within a temperature range of 450°C to 600°C after hot-dip galvanizing. When alloying is performed at temperatures exceeding 600°C, the untransformed austenite transforms into pearlite, failing to ensure the desired residual austenite area ratio, and sometimes reducing ductility. Therefore, it is preferable to perform the alloying process within a temperature range of 450°C to 600°C.

[0134] There are no particular limitations on the conditions of other manufacturing methods. From a productivity point of view, the above-mentioned annealing is preferably carried out by a continuous annealing equipment. In addition, a series of treatments such as annealing, galvanizing, and alloying of the galvanized layer are preferably carried out by a hot-dip galvanizing line (CGL, Continuous Galvanizing Line).

[0135] It should be noted that for purposes such as shape correction and surface roughness adjustment, the aforementioned "high-strength steel plates" and "high-strength hot-dip galvanized steel plates" can undergo surface finishing rolling. The reduction rate of surface finishing rolling is preferably in the range of 0.1% to 2.0%. Below 0.1%, the effect is small and difficult to control; therefore, this is considered the lower limit of the good range. Furthermore, above 2.0%, productivity decreases significantly; therefore, this is considered the upper limit of the good range. It should be noted that surface finishing rolling can be performed online or offline. Furthermore, surface finishing rolling with the target reduction rate can be performed in one go or in multiple stages. Additionally, various coating treatments such as resin or grease coating can be applied.

[0136] Example

[0137] Steel with the composition shown in Table 1, with the balance consisting of Fe and unavoidable impurities, was smelted in a converter and produced into billets by continuous casting. The resulting billets were then reheated to 1250°C and subjected to the conditions shown in Tables 2 and 3 to obtain high-strength cold-rolled steel sheets (CR). Further galvanizing was then performed to obtain hot-dip galvanized steel sheets (GI) and alloyed hot-dip galvanized steel sheets (GA). It should be noted that the thickness of CR, GI, and GA sheets is 1.0 mm or more and 1.8 mm or less. Regarding the hot-dip galvanizing bath, a zinc bath containing 0.19% by mass Al was used for hot-dip galvanized steel sheets (GI), and a zinc bath containing 0.14% by mass Al was used for alloyed hot-dip galvanized steel sheets (GA). The bath temperature was set at 465°C. The coating adhesion was set at 45 g / m² per single side. 2 (Double-sided coating), GA is adjusted to ensure that the Fe concentration in the coating is 9% by mass or more and 12% by mass or less. The steel microstructure of the obtained steel plate cross section is observed using the above method, and the tensile properties, hole expansion properties, bending properties, and coating properties are investigated. The results are shown in Tables 4 to 6.

[0138]

[0139] The martensitic phase transformation start temperature, as well as the Ac1 and Ac3 phase transformation points, are calculated using the following formula.

[0140] Martensitic transformation onset temperature (°C) = 550 - 350 × (%C) - 40 × (%Mn) - 10 × (%Cu) - 17 × (%Ni) - 20 × (%Cr) - 10 × (%Mo) - 35 × (%V) - 5 × (%W) + 30 × (%Al)

[0141] Ac1 phase transition point (°C) = 751 - 16 × (% C) + 11 × (% Si) - 28 × (% Mn) - 5.5 × (% Cu) - 16 × (% Ni) + 13 × (% Cr) + 3.4 × (% Mo)

[0142] Ac3 phase transition point (°C) = 910 - 203√(%C) + 45×(%Si) - 30×(%Mn) - 20×(%Cu) - 15×(%Ni) + 11×(%Cr) + 32×(%Mo) + 104×(%V) + 400×(%Ti) + 200×(%Al)

[0143] Here, (%C), (%Si), (%Mn), (%Ni), (%Cu), (%Cr), (%Mo), (%V), (%Ti), (%W), and (%Al) represent the content (mass%) of each element, which is set to zero if the element is not present.

[0144]

[0145]

[0146] [Table 4]

[0147]

[0148] Underlined portion: indicates outside the scope of this invention.

[0149] F: Ferrite, M: Fresh martensite, RA: Retained austenite

[0150] TM: Tempered martensite, B: Bainite

[0151] [Table 5]

[0152]

[0153] Underlined portion: indicates outside the scope of this invention.

[0154] F: Ferrite, RA: Retained austenite, P: Pearlite, θ: Carbides (cementite, etc.)

[0155] [Table 6]

[0156]

[0157] Underlined portion: indicates outside the scope of this invention.

[0158] In the tensile test, JIS 5 test pieces were used, cut with the tensile direction perpendicular to the rolling direction of the steel sheet. The test was conducted according to JIS Z 2241 (2011), measuring TS (tensile strength) and EL (total elongation). In the case of coated steel sheets, the post-coating ductility (EL / EL') was also measured. Here, EL' represents the total elongation without immersion in a coating bath during the rolling process; for cold-rolled steel sheets, EL = EL'. Furthermore, regarding mechanical properties, the following conditions were considered good.

[0159] For pressures above 980 MPa and below 1180 MPa, EL ≥ 20% and EL / EL' ≥ 0.7.

[0160] For pressures above 1180 MPa, EL ≥ 12% and EL / EL' ≥ 0.7.

[0161] Hole expansion performance was assessed according to JIS Z 2256 (2010). Each steel plate was cut into 100mm x 100mm pieces, and holes with a diameter of 10mm were punched with a gap of 12% ± 1%. Then, using a die with an inner diameter of 75mm and a blank holder force of 9 tons, a 60° conical punch was pressed into the hole. The hole diameter at the point where cracking occurred was measured, and the limiting expansion rate λ (%) was calculated using the following formula. Hole expansion performance was evaluated based on this limiting expansion rate.

[0162] Limiting porosity λ (%) = {(D f -D0) / D0}×100

[0163] Among them, D f D0 is the pore diameter (mm) at which cracks occur, and D0 is the initial pore diameter (mm). It should be noted that in this invention, for each TS range, the following conditions are considered good.

[0164] For pressures above 980 MPa but below 1180 MPa, λ ≥ 15%.

[0165] For pressures above 1180 MPa, λ ≥ 25%.

[0166] Regarding the bending test, bending test pieces with a width of 30 mm and a length of 100 mm were cut from each annealed steel sheet with the rolling direction as the bending axis. The test was conducted based on the V-block method according to JIS Z 2248 (1996). Tests were performed at a pressing speed of 100 mm / s for n=3 tests at each bending radius. For the outer side of the bent portion, the presence or absence of cracks was determined using a stereomicroscope. The smallest bending radius without cracking was taken as the ultimate bending radius R (mm). It should be noted that in this invention, the condition satisfying the ultimate bending radius R / t ≤ 2.5 (t: steel sheet thickness) under 90° V bending is considered to indicate good bending performance of the steel sheet.

[0167] The plating performance is evaluated by appearance. A case where there are no defects such as no plating, uneven alloying, or other defects that impair surface quality, and where appropriate surface quality is ensured, is categorized as ○. A case where there is no uneven color and other defects, resulting in an excellent appearance, is categorized as ◎. A case where some minor defects are observed is categorized as △. A case where a large number of surface defects are observed is categorized as ×. Cases ◎, ○, and △ are considered within the scope of this invention.

[0168] The high-strength steel sheets of the present invention all have a strength (TS) of 980 MPa or higher and excellent formability. On the other hand, in the comparative examples, at least one of the following characteristics is poor: TS, EL, post-plating ductility, λ, bendability, and plating properties.

[0169] Industrial availability

[0170] According to the present invention, a high-strength steel sheet with excellent formability and a tensile strength (TS) of 980 MPa or higher can be obtained. By applying the high-strength steel sheet of the present invention to, for example, automotive structural components, fuel efficiency improvements resulting from vehicle body lightweighting can be achieved, demonstrating significant industrial application value.

Claims

1. A high-strength steel plate, comprising: It contains, by mass%, C: 0.030% or more and 0.250% or less, Si: 0.01% or more and 3.00% or less, Mn: 2.00% or more and 8.00% or less, P: 0.100% or less, S: 0.0200% or less, N: 0.0100% or less, Al: 0.001% or more and 2.000% or less, with the balance being Fe and unavoidable impurities; and Steel microstructure with an area ratio of 1% to 40% ferrite, 1% to 20% fresh martensite, 35% to 90% bainite and tempered martensite, and 6% or more retained austenite. The value obtained by dividing the average Mn content (by mass%) in retained austenite by the average Mn content (by mass%) in ferrite is 1.1 or more, and the value obtained by dividing the average C content (by mass%) in retained austenite with an aspect ratio of 2.0 or more by the average C content (by mass%) in ferrite is 3.0 or more. The value obtained by dividing the amount of carbon in all the retained austenite by the amount of carbon in the T0 composition is 1.0 or higher. The T0 composition is the composition in which the free energies of fcc and bcc are equal at the reheating temperature before immersion in the zinc bath, i.e., 120 to 480°C.

2. The high-strength steel plate according to claim 1, wherein, The composition, by mass%, further contains at least one element selected from the following: Ti: less than 0.200%, Nb: less than 0.200%, V: less than 0.500%, W: less than 0.500%, B: less than 0.0050%, Ni: less than 1.000%, Cr: less than 1.000%, Mo: less than 1.000%, Cu: less than 1.000%, Sn: less than 0.200%, Sb: less than 0.200%, Ta: less than 0.100%, Zr: less than 0.200%, Ca: less than 0.0050%, Mg: less than 0.0050%, and REM: less than 0.0050%.

3. The high-strength steel plate according to claim 1 or 2, wherein, The value obtained by dividing the area ratio of blocky retained austenite by the area ratio of all retained austenite and blocky fresh martensite is less than 0.

5.

4. The high-strength steel plate according to claim 1 or 2, further comprising a galvanized layer on its surface.

5. The high-strength steel plate according to claim 3, further comprising a galvanized layer on its surface.

6. The high-strength steel plate according to claim 4, wherein, The zinc plating layer is an alloyed zinc plating layer.

7. The high-strength steel plate according to claim 5, wherein, The zinc plating layer is an alloyed zinc plating layer.

8. A method for manufacturing a high-strength steel plate, which is the method for manufacturing a high-strength steel plate according to any one of claims 1 to 3, wherein, A steel billet having the composition described in claim 1 or 2 is heated, hot-rolled at a temperature of 750°C or higher and 1000°C or lower at the finishing mill exit, coiled and cold-rolled at 300°C or higher and 750°C or lower, then held at a temperature range of -50°C or higher from the Ac3 phase transformation point for 20 seconds or more and 1800 seconds or less, cooled to a cooling stop temperature below the martensitic phase transformation start temperature, reheated to a reheating temperature range of 120°C or higher and 450°C or lower, held at the reheating temperature for 2 seconds or more and 1800 seconds or less, and cooled to room temperature. Then, held at a temperature range of -20°C or higher from the Ac1 phase transformation point for 20 seconds or more and 600 seconds or less, cooled to a cooling stop temperature below the martensitic phase transformation start temperature, reheated to a reheating temperature range of 120°C or higher and 480°C or lower, held at the reheating temperature for 2 seconds or more and 600 seconds or less, and then cooled to room temperature.

9. The method for manufacturing high-strength steel plate according to claim 8, wherein, Further galvanizing treatment is carried out.

10. The method for manufacturing high-strength steel plate according to claim 9, wherein, Following the galvanizing process, an alloying process is performed at a temperature above 450°C and below 600°C.

11. A method for manufacturing a high-strength steel plate according to any one of claims 8 to 10, wherein, After coiling and before cold rolling, the temperature is maintained for more than 1800 s in a temperature range below the Ac1 phase transformation point.

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