High-strength steel sheet and method for manufacturing the same

Abstract

The present invention aims to provide a high-strength steel sheet having a TS of 980 MPa or more and excellent ductility, hole expandability, bendability, and hydrogen embrittlement bendability, and a method for manufacturing the same. A high-strength steel sheet having a prescribed composition, a steel structure in which ferrite is 1% or more and 40% or less in area ratio, fresh martensite is less than 1.0%, the sum of bainite and tempered martensite is 40% or more and 90% or less, retained austenite is 6% or more, a value obtained by dividing the average Mn amount (mass%) in the retained austenite by the average Mn amount (mass%) in the ferrite is 1.1 or more, and a value obtained by dividing the average C amount (mass%) in the retained austenite having a length-diameter ratio of 2.0 or more by the average C amount (mass%) in the ferrite is 3.0 or more, and the diffusible hydrogen amount in the steel is 0.3 mass ppm or less.

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 engineering, and methods for manufacturing them. In particular, it is desirable to obtain high-strength steel sheets with a tensile strength (TS) of 980 MPa or higher, low hydrogen content within the steel, and excellent resistance to hydrogen embrittlement during bending. 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 reducing the weight of the vehicle body by increasing the strength of body materials to achieve thinner walls. On the other hand, increasing the strength of steel sheets leads to a decrease in formability. Furthermore, annealing in a reducing atmosphere containing hydrogen allows hydrogen to penetrate into the steel sheet, reducing its bendability and other formability properties. Therefore, there is a need to develop materials that possess high strength and high formability, while also exhibiting resistance to hydrogen embrittlement.

[0003] As a high-strength and highly ductile steel sheet, 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 forming. 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] Furthermore, Patent Document 2 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.

[0006] Furthermore, Patent Document 3 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 dual-phase region of ferrite and austenite. As a result, a large aspect ratio retained austenite 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.

[0007] Furthermore, Patent Document 4 discloses a method in which annealed steel sheets, hot-dip galvanized steel sheets, or alloyed hot-dip galvanized steel sheets are held at a temperature range of 50°C to 300°C for 1800 seconds to 43200 seconds, thereby reducing the hydrogen content in the steel. However, no research was conducted on the improvement in bendability resulting from the reduction in hydrogen content in the steel.

[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. 2003-138345

[0012] Patent Document 3: Patent No. 6123966

[0013] Patent Document 4: International Publication No. 2019 / 188642 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, excellent formability, and low hydrogen content within the steel, as well as a method for manufacturing the same. Formability, as described herein, refers to ductility, pore-expanding ability, and bending properties.

[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 aforementioned cooling, the steel is held at a temperature range of -20°C to the Ac1 phase transformation point for 20 seconds to 600 seconds, then cooled to a cooling stop temperature below the martensitic transformation start temperature, and reheated to a reheating temperature range of 120°C to 480°C. Then, it is held at the aforementioned reheating temperature for 2 seconds to 600 seconds, followed by plating treatment as needed, and then cooled to a temperature range of room temperature to below the martensitic transformation start temperature. Then, it is further held at a temperature range of 50°C to 400°C for 2 seconds to 2 seconds, thereby efficiently removing hydrogen and improving resistance to hydrogen embrittlement during bending. The steel sheet manufactured in the above manner has a steel microstructure with ferrite content of 1% to 40% and fresh martensite content (by area ratio), less than 1.0% and the sum of bainite and tempered martensite content of 40% to 90% and retained austenite content of 6% or more. Furthermore, it is known that a high-strength steel sheet with excellent formability and resistance to hydrogen embrittlement bending can be manufactured. The steel sheet is characterized in that, in the steel microstructure, 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, and the diffusible hydrogen content in the steel is 0.3 ppm by mass or less.

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

[0021] This invention is based on the above insights, and its main points are as follows.

[0022] [1] A high-strength steel plate comprising, 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 ferrite comprising 1% or more and 40% by area. The steel microstructure has the following characteristics: less than 6% of fresh martensite, less than 1.0% of fresh martensite, a sum of bainite and tempered martensite of 40% to 90% or more, and a retained austenite of 6% or more. The average Mn content (mass%) in the retained austenite divided by the average Mn content (mass%) in the ferrite has a value of 1.1 or more, and the average C content (mass%) in the retained austenite divided by the average C content (mass%) in the ferrite has a length-to-diameter ratio of 2.0 or more has a value of 3.0 or more. The diffusible hydrogen content in the steel is less than 0.3 ppm by mass.

[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 sheet according to any one of [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 finishing mill exit side, coiled and cold-rolled at 300°C or higher and 750°C or lower, and then held at a temperature range of -50°C or higher from the Ac3 phase transformation point for 20s or more and 1800s or less, cooled to a cooling stop temperature below the martensitic phase transformation start temperature, and then heated to 120°C or higher and 450°C or lower. After reheating to a temperature within the range, hold at the reheating temperature for 2 seconds to 1800 seconds, then cool to room temperature. Then, hold at a temperature range above -20°C from the Ac1 phase transformation point for 20 seconds to 600 seconds, then cool to a cooling stop temperature below the martensitic phase transformation start temperature. After reheating to a temperature range above 120°C to 480°C, hold at the reheating temperature for 2 seconds to 600 seconds, then cool to a temperature range above room temperature and below the martensitic phase transformation start temperature, and further hold at a temperature range above 50°C to 400°C for 2 seconds or more.

[0028] [7] In the method for manufacturing high-strength steel plate according to [6], after reheating to a reheating temperature in the range of 120°C or higher and 480°C or lower, after holding at the reheating temperature for 2 seconds or more and 600 seconds or less, and before cooling to room temperature or higher and below the martensitic transformation start temperature, a plating process is further performed.

[0029] [8] The method for manufacturing high-strength steel plate according to [7], wherein galvanizing is performed in the above-mentioned plating process.

[0030] [9] According to the manufacturing method of high-strength steel plate described in [8], alloying treatment is carried out at a temperature above 450°C and below 600°C after the above-mentioned galvanizing treatment.

[0031]

[10] The method for manufacturing high-strength steel sheet according to any one of [6] to [9], 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.

[0032] Invention Effects

[0033] According to the present invention, a high-strength steel sheet with a tensile strength (TS) of over 980 MPa, excellent formability after plating treatment, and particularly excellent ductility, hole expansion, and bending properties 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 application value. Detailed Implementation

[0034] 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.

[0035] (1) The reasons for limiting the composition of steel to the above range in this invention will be explained.

[0036] C: Above 0.030% and below 0.250%

[0037] Carbon (C) is an element required to increase strength by inducing low-temperature phase transformations such as martensite. Additionally, 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 formation occurs, resulting in unsatisfactory strength.

[0038] Furthermore, it is difficult to ensure a sufficient residual austenite area ratio, 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 expansion tests, micro-voids at the martensite grain boundaries increase, and crack propagation occurs, reducing porosity. Additionally, the hardening of the weld and heat-affected zone is significant, reducing the mechanical properties of the weld and thus deteriorating spot weldability, arc weldability, etc. 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.

[0039] Si: ≥0.01% and ≤3.00%

[0040] 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 Si content exceeding 3.00% causes embrittlement of the steel, making it difficult to ensure ductility. Furthermore, it deteriorates surface properties due to the formation of red oxide scale, etc., and also leads to a reduction in coating quality. Therefore, Si content 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%.

[0041] Mn: ≥2.00% and ≤8.00%

[0042] Mn is an extremely important element in this invention. Mn stabilizes retained austenite, effectively 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, an excess of Mn exceeding 8.00% can lead to the formation of an inhomogeneous banded structure due to Mn segregation, deteriorating bending properties. 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.

[0043] P: below 0.100%

[0044] P is an element that provides solid solution strengthening and can be included 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. Furthermore, a preferred upper limit is set to 0.050% or less.

[0045] S: below 0.0200%

[0046] S segregation at grain boundaries 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 a production cost perspective, 0.0001% or more is preferred. Therefore, the S amount is set to 0.0200% or less. The upper limit is preferably set to 0.0100% or less, and more preferably 0.0050% or less.

[0047] N: below 0.0100%

[0048] Nitrogen (N) is an element that degrades the aging resistance of steel. The degradation of aging resistance becomes particularly 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 higher. Therefore, the N content is set to 0.0100% or lower. A more preferred lower limit is set to 0.0010% or higher. A more preferred upper limit is set to 0.0070% or lower.

[0049] Al: Above 0.001% and below 2.000%

[0050] Al is an element that expands the two-phase region between ferrite and austenite, reduces the dependence of mechanical properties on annealing temperature, and thus contributes to 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, effectively improving the cleanliness of steel, and is preferably added during the deoxidation process. However, a high content 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 0.200% or more. A preferred upper limit is 1.200% or less.

[0051] 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.

[0052] Ti: below 0.200%

[0053] 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 Ti is included, its content is set to 0.200% or less. The preferred lower limit is set to 0.005% or more, more preferably 0.010% or more. The preferred upper limit is set to 0.100% or less.

[0054] Nb: below 0.200%, V: below 0.500%, W: below 0.500%

[0055] Nb, V, and W are effective in precipitation strengthening of steel, similar to the effect of Ti. 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, the micropores at the martensite grain boundaries increase, and crack propagation occurs, sometimes reducing porosity. Therefore, when Nb is included, its content is set to 0.200% or less, preferably a lower limit of 0.005% or more, more preferably 0.010% or more, and preferably an upper limit of 0.100% or less. When V and W are included, their contents are each set to 0.500% or less, preferably a lower limit of 0.005% or more, more preferably 0.010% or more. The preferred upper limit values ​​are each set to below 0.300%.

[0056] B: Below 0.0050%

[0057] 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 B is included, its content is set to 0.0050% or less. The preferred lower limit is set to 0.0003% or more, more preferably 0.0005% or more. Furthermore, the preferred upper limit is set to 0.0030% or less.

[0058] Ni: below 1.000%

[0059] 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 it is included at a concentration exceeding 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 Ni is included, its content is set to 1.000% or less, preferably 0.005% or more and 1.000% or less.

[0060] Cr: less than 1.000%, Mo: less than 1.000%

[0061] Cr and Mo have the effect of improving the balance between strength and ductility, and therefore can be included as needed. However, when Cr and Mo are contained 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 these elements are included, 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%.

[0062] Cu: below 1.000%

[0063] 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 it contains more than 1.000%, the area ratio of hard martensite becomes too large, and during the porosity expansion test, the micro-voids at the grain boundaries of the martensite increase, and crack propagation occurs, reducing porosity. Therefore, when Cu is present, its amount is set to 1.000% or less, preferably 0.005% or more and 1.000% or less.

[0064] Sn: less than 0.200%, Sb: less than 0.200%

[0065] 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 included 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 content of any one of these elements exceeding 0.200% will lead to a decrease in toughness. Therefore, when Sn and Sb are included, their contents are each set to 0.200% or less, preferably 0.002% or more and 0.200% or less.

[0066] Ta: below 0.100%

[0067] Like Ti and Nb, Ta forms alloy carbides and alloy carbonitrides, contributing to increased strength. 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 included, its content is set to 0.100% or less, preferably 0.001% or more and 0.100% or less.

[0068] Zr: below 0.200%

[0069] 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 Zr content exceeding 0.200% can lead to an increase in inclusions and other defects, causing surface and internal defects. Therefore, when Zr is present, its content is set to 0.200% or less, preferably 0.0005% or more and 0.0050% or less.

[0070] Ca: less than 0.0050%, Mg: less than 0.0050%, REM: less than 0.0050%

[0071] 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 amounts of each exceeding 0.0050% can lead to an increase in inclusions and other defects, causing surface and internal defects. Therefore, when Ca, Mg, and REM are present, their contents are each set to 0.0050% or less, preferably 0.0005% or more and 0.0050% or less.

[0072] The balance other than the above components is Fe and unavoidable impurities.

[0073] (2) Next, the steel structure will be described.

[0074] Ferrite area fraction: 1% or more and 40% or less

[0075] 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.

[0076] Area fraction of fresh martensite: less than 1.0%

[0077] The significant hardness difference between fresh martensite and the soft ferrite phase degrades porosity during punching. Therefore, to ensure good porosity, the area fraction of fresh martensite needs to be less than 1.0%.

[0078] The sum of the area ratios of bainite and tempered martensite is 40%–90%.

[0079] 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 40%, good porosity is not achieved. Therefore, the sum of the area ratios of bainite and tempered martensite needs to be 40% 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 50% or more and 85% or less.

[0080] 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), 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.

[0081] Area ratio of retained austenite: 6% or more

[0082] 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.

[0083] 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 the 1 / 4 position of the plate thickness, another 0.1 mm was further ground away by chemical grinding. For the ground surface at the 1 / 4 position of the plate thickness, 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.

[0084] 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.

[0085] 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.

[0086] 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.

[0087] 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.

[0088] The amounts of C and Mn in retained austenite and ferrite can be quantified using FE-EPMA (Field Emission-Electron Probe Micro Analyzer) to measure the distribution of Mn in each phase of the rolled section at 1 / 4 of the plate thickness. The Mn content is calculated by averaging the results of quantitative analysis of 30 retained austenite grains and 30 ferrite grains.

[0089] 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.

[0090] The diffusible hydrogen content in the steel is below 0.3 ppm by mass.

[0091] To ensure good resistance to hydrogen embrittlement during bending, it is important to set the diffusible hydrogen content in the steel to below 0.3 ppm by mass, preferably below 0.20 ppm by mass. It should be noted that there is no specific lower limit for the diffusible hydrogen content in the steel; considering production technology constraints, the diffusible hydrogen content in the steel can be above 0.01 ppm by mass.

[0092] The method for determining the diffusible hydrogen content in steel is described below. A test piece with a length of 30 mm and a width of 5 mm is cut from the product coil. In the case of hot-dip galvanized steel sheet or alloyed hot-dip galvanized steel sheet, the hot-dip galvanized layer or alloyed hot-dip galvanized layer of the test piece is removed by grinding or alkali. Then, the amount of hydrogen released from the test piece is determined by thermal desorption spectrometry (TDS). Specifically, the test piece is continuously heated from room temperature to 300°C at a heating rate of 200°C / hour, and then cooled to room temperature. The cumulative amount of hydrogen released from the test piece from room temperature to 210°C is measured as the diffusible hydrogen content in the steel.

[0093] The value obtained by dividing the area ratio of massive retained austenite by the area ratio of all retained austenite and massive fresh martensite is less than 0.5.

[0094] Due to the constraint from surrounding grains, bulk retained austenite exhibits high stability. Therefore, during punching, martensitic transformation occurs in high-strain regions, increasing the hardness difference with surrounding grains and sometimes deteriorating porosity. Thus, it is preferable that the area ratio of bulk retained austenite divided by the area ratio of all retained austenite and bulk fresh martensite is 0.5 or less. More preferably, it is 0.4 or less. It should be noted that bulk retained austenite refers to austenite with an aspect ratio less than 2.0. There is no limitation on the average grain size of bulk retained austenite, but an average grain size of, for example, 3 μm or less can be considered. This average grain size can be determined using conventionally known methods, such as image analysis of microstructure images of bulk retained austenite obtained using a scanning electron microscope (SEM).

[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] 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.

[0097] 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.

[0098] (3) Next, the manufacturing conditions will be explained.

[0099] Heating temperature of steel billet

[0100] 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 precipitate 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 sometimes 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.

[0101] 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.

[0102] Finishing mill exit temperature of hot-rolled steel: above 750℃ and below 1000℃

[0103] 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.

[0104] Coiling temperature after hot rolling: above 300℃ and below 750℃

[0105] 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.

[0106] 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.

[0107] 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.

[0108] cold rolling

[0109] 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%.

[0110] It remains below the Ac1 phase transition point for more than 1800 seconds.

[0111] Holding the steel sheet at temperatures below the Ac1 transformation point for more than 1800 seconds allows for subsequent cold-rolling softening, and this process can be performed as needed. When holding at temperatures above the Ac1 transformation point, Mn enriches in the austenite, and upon cooling, hard martensite and retained austenite are formed, sometimes preventing the steel sheet from softening. Furthermore, holding for less than 1800 seconds fails to remove the strain from hot rolling, sometimes also preventing the steel sheet from softening.

[0112] 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.

[0113] 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).

[0114] When held at temperatures below the Ac3 transformation point (-50°C), Mn enriches in austenite and does not undergo martensitic transformation during cooling, failing to produce nuclei of retained austenite with a large aspect ratio. As a result, in subsequent annealing processes (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, resulting in an undesirable microstructure. Holding for less than 20 seconds does not allow for sufficient recrystallization, resulting in an undesirable microstructure and reduced porosity. Furthermore, sufficient Mn surface enrichment to ensure subsequent coating quality is not achieved. On the other hand, holding for more than 1800 seconds leads to excessive Mn surface enrichment, deteriorating coating quality. Moreover, the austenite grains coarsen during annealing, resulting in the retention of nuclei of retained austenite with a small aspect ratio during subsequent cooling, leading to an undesirable microstructure and reduced ductility, porosity, and bendability.

[0115] Cooling to the cooling stop temperature below the martensitic transformation start temperature

[0116] When the cooling stop temperature exceeds the martensitic transformation initiation 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. Preferably, the martensitic transformation initiation temperature is above -250°C and below -50°C.

[0117] After reheating to a reheating temperature within the range of 120°C to 450°C, hold at the above reheating temperature for 2 seconds to 1800 seconds, and then cool to room temperature.

[0118] At reheating temperatures below 120°C, carbon will not accumulate in the retained austenite formed during subsequent annealing processes, resulting in an undesirable microstructure. 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 failing to yield the desired microstructure. Similarly, holding the temperature for less than 2 seconds also fails to produce nuclei of retained austenite with large aspect ratios, resulting in an undesirable microstructure.

[0119] Furthermore, when the temperature exceeds 1800s, the nuclei of the retained austenite with a large aspect ratio decompose, while the retained austenite with a small aspect ratio increases. Mn will not be enriched in the retained austenite, and the desired microstructure cannot be obtained.

[0120] 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.

[0121] 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).

[0122] Maintaining the temperature within a range of 20°C above the Ac1 transformation point for 20 seconds to 600 seconds is an extremely important inventive requirement in this invention. When the temperature range is below the Ac1 transformation point of 20°C and the holding time is less than 20 seconds, the amount of austenite in the annealing process is low, the area ratio of ferrite increases, and it is difficult to ensure the total ductility (TS). Furthermore, the carbides formed during heating are not completely dissolved, making it difficult to ensure a sufficient area ratio of retained austenite, thus reducing ductility.

[0123] Preferably, the temperature is above the Ac1 phase transformation point. More preferably, it is above the Ac1 phase transformation point +20°C and below the Ac3 phase transformation point +50°C. Furthermore, when held for more than 600 s, the austenite coarsens during annealing, and therefore, the diffusion of Mn into the austenite becomes insufficient, and it cannot be enriched, so that a sufficient area ratio of residual austenite can not be obtained to ensure ductility.

[0124] Cooling to the cooling stop temperature below the martensitic transformation start temperature

[0125] 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.

[0126] After reheating to a reheating temperature within the range of 120°C to 480°C, maintain the reheating temperature at the above-mentioned reheating temperature for 2 seconds to 600 seconds.

[0127] 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, leading to an undesirable microstructure. Furthermore, carbide precipitation and reduced austenite stabilization result in an undesirable amount of retained austenite.

[0128] Furthermore, when the holding time is less than 2 seconds, not only is the fresh martensite not tempered, but carbon also does not accumulate in the γ with a large aspect ratio, resulting in an undesirable microstructure. On the other hand, when the holding time exceeds 600 seconds, carbides precipitate during the bainitic transformation, reducing the carbon content in the retained austenite and also resulting in an undesirable microstructure.

[0129] Plating treatment

[0130] The obtained high-strength steel sheet is then coated 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.

[0131] 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.

[0132] Cooling stop temperature above room temperature and below the martensitic phase transformation start temperature

[0133] When cooling stops above the martensitic transformation temperature, the amount of slowly diffusing austenite increases during subsequent reheating, and the amount of diffusible hydrogen in the steel is not sufficiently reduced. Therefore, cooling to below the martensitic transformation initiation temperature is necessary. Preferably, the temperature is above 50°C and below the martensitic transformation initiation temperature -30°C.

[0134] Maintain for more than 2 seconds within a temperature range above 50℃ and below 400℃.

[0135] As a final heat treatment, holding the steel at a temperature range of 50°C to 400°C for at least 2 seconds is an important inventive element in this invention. When held at temperatures below 50°C or for less than 2 seconds, an excessive amount of fresh martensite is generated, and diffusible hydrogen in the steel is not released from the steel sheet, thus reducing its resistance to hydrogen embrittlement during bending. On the other hand, when held at temperatures exceeding 400°C, insufficient volume fraction of retained austenite is obtained due to the decomposition of the retained austenite, reducing the ductility of the steel. There is no specific upper limit to the holding time; considering production technology constraints, it can be below 43200 seconds.

[0136] 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, hot-dip galvanizing, and alloying of the galvanized layer are preferably carried out by a hot-dip galvanizing line (CGL, Continuous Galvanizing Line).

[0137] 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.

[0138] Example

[0139] 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, and bending properties are investigated. The results are shown in Tables 4 to 6.

[0140]

[0141] The martensitic phase transformation start temperature, as well as the Ac1 and Ac3 phase transformation points, are calculated using the following formula.

[0142] Martensitic transformation onset temperature (°C) = 550 - 350 × (%C) - 40 × (%Mn) - 10 × (%Cu) - 17 × (%Ni) - 20 × (%Cr) - 10 × (%Mo) - 35 × (%V) - 5 × (%W) + 30 × (%Al)

[0143] Ac1 phase transition point (°C) = 751 - 16 × (% C) + 11 × (% Si) - 28 × (% Mn) - 5.5 × (% Cu) - 16 × (% Ni) + 13 × (% Cr) + 3.4 × (% Mo)

[0144] 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)

[0145] 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.

[0146]

[0147]

[0148] [Table 4]

[0149]

[0150] Underlined portion: indicates outside the scope of this invention.

[0151] F: Ferrite, M: Fresh martensite, RA: Retained austenite

[0152] [Table 5]

[0153]

[0154] Underlined portion: indicates outside the scope of this invention. 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 plate. The test was conducted according to JIS Z 2241 (2011), and TS (tensile strength) and EL (total elongation) were measured. Regarding mechanical properties, the following conditions were considered good.

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

[0160] TS: When the pressure is above 1180MPa, EL ≥ 12%.

[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 Z2248 (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. 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 properties of the steel sheet.

[0167] Regarding the hydrogen embrittlement resistance bending characteristics, the bending test described above is evaluated as follows. If the value obtained by dividing the measured R / t of the steel plate by (R / t)' when the hydrogen content in the same steel plate is 0.00 ppm by mass is less than 1.4, it is considered to have good hydrogen embrittlement resistance in this invention. It should be noted that (R / t)' is determined by the following method: the hydrogen content in the steel plate is reduced by placing it in the atmosphere for a long time, and then the hydrogen content in the steel is confirmed to reach 0.00 ppm by mass using TDS (Thermal Desorption Spectrometry) before a bending test is performed.

[0168] The high-strength steel sheets of the present invention all exhibit 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—TS, EL, λ, bendability, and resistance to hydrogen embrittlement during bending—is poor.

[0169] Industrial availability

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

Claims

1. A high-strength steel sheet having: a composition consisting of, in 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, and the balance consisting of Fe and unavoidable impurities; and a steel structure in which ferrite is 1% or more and 40% or less in area ratio, fresh martensite is less than 1.0%, the sum of bainite and tempered martensite is 40% or more and 90% or less, and retained austenite is 6% or more, a value obtained by dividing an average Mn amount in mass% in the retained austenite by an average Mn amount in mass% in the ferrite is 1.1 or more, and a value obtained by dividing an average C amount in mass% in the retained austenite in which the aspect ratio is 2.0 or more by an average C amount in mass% in the ferrite is 3.0 or more, an amount of diffusible hydrogen in the steel is 0.3 mass ppm or less. the composition further contains, in mass%, at least one element selected from the group consisting of Ti: 0.200% or less, Nb: 0.200% or less, V: 0.500% 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: 0.100% or less, Zr: 0.200% or less, Ca: 0.0050% or less, Mg: 0.0050% or less, REM: 0.0050% or less. a value obtained by dividing an area ratio of blocky retained austenite in which the aspect ratio is less than 2.0 by the sum of area ratios of the total retained austenite and fresh martensite is 0.5 or less.

2. The high-strength steel sheet according to claim 1, wherein 4. The high-strength steel sheet according to claim 1 or 2, further having a galvanized layer on a surface.

3. The high-strength steel sheet according to claim 1 or 2, wherein, 5. The high-strength steel sheet according to claim 3, further having a galvanized layer on a surface. the galvanized layer is an alloyed galvanized layer. the galvanized layer is an alloyed galvanized layer.

6. The high-strength steel sheet according to claim 4, wherein 8. A method of manufacturing a high-strength steel sheet according to any one of claims 1 to 3, wherein 7. The high-strength steel sheet according to claim 5, wherein ​ ​ The steel slab having the component 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 exit side of finish rolling, coiled at a temperature of 300°C or higher and 750°C or lower, cold-rolled, then held at a temperature of Ac3 transformation point - 50°C or higher for 20 seconds or more and 1800 seconds or less, cooled to a cooling stop temperature of a temperature below the martensite transformation start temperature, reheated to a reheating temperature in a 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, then cooled to room temperature, held at a temperature of Ac1 transformation point - 20°C or higher for 20 seconds or more and 600 seconds or less, cooled to a cooling stop temperature of a temperature below the martensite transformation start temperature, reheated to a reheating temperature in a 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, then cooled to a temperature of room temperature or higher and below the martensite transformation start temperature, and held at a temperature in a range of 50°C or higher and 400°C or lower for 2 seconds or more.

9. The method of producing a high-strength steel sheet according to claim 8, wherein After the reheating to the reheating temperature in the range of 120°C or higher and 480°C or lower and before the cooling to the temperature of room temperature or higher and below the martensite transformation start temperature, plating treatment is further performed.

10. The method of producing a high-strength steel sheet according to claim 9, wherein In the plating treatment, galvanizing treatment is performed.

11. The method of producing a high-strength steel sheet according to claim 10, wherein After the galvanizing treatment, alloying treatment is performed at a temperature of 450°C or higher and 600°C or lower.

12. The method of producing a high-strength steel sheet according to any one of claims 8 to 11, wherein After the coiling and before the cold-rolling, the steel is held at a temperature below the Ac1 transformation point for more than 1800 seconds. In the plating treatment, galvanizing treatment is performed.

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