HIGH STRENGTH STEEL SHEET AND METHOD FOR MANUFACTURING IT.

Abstract

One objective is to provide a high-strength steel sheet having a tensile strength of 980 MPa or more and excellent formability, and a method for manufacturing the high-strength steel sheet. A high-strength steel sheet with a predetermined chemical composition is provided and manufactured under optimum conditions. The high-strength steel sheet has a steel microstructure that includes, by area, ferrite: 30% or more and 80% or less, martensite: 5% or more and 35% or less, and retained austenite: 8% or more, wherein the ratio of the grain area fraction of retained austenite, grains having an aspect ratio of 2.0 or more and a minor axis length of 1µm or less, divided by the total area fraction of retained austenite is 0.3 or more, wherein the ratio of the average Mn content (% by mass) in the retained austenite divided by the average Mn content (% by mass) in the ferrite is 1.5 or more, and the product of the ratio of the average Mn content (% by mass) in the retained austenite divided by the average Mn content (% by mass) in the ferrite and the average aspect ratio of the retained austenite is 3.0 or more, the ratio of the average C content (% by mass) in the retained austenite divided by the average C content (% by mass) in the ferrite is 3.0 or more, and wherein the ratio of the average C content (% by mass) in the retained austenite divided by the average Mn content (% by mass) in the retained austenite is 0.05 or more.

Description

HIGH-STRENGTH STEEL SHEET AND METHOD FOR MANUFACTURING IT Technical Field The present invention relates to a high-strength steel sheet with excellent formability suitable for use as a component in various industries, such as the automotive and electrical industries, and a method for manufacturing the high-strength steel sheet. Specifically, the present invention aims to produce a high-strength steel sheet having a tensile strength (TS) of 980 MPa or higher and exhibiting excellent ductility, flanging, and bending capabilities. Background Improving fuel economy in automobiles has been a major concern from a global environmental perspective. Consequently, efforts have been made to increase the strength of automotive body materials to enable reductions in body thickness and weight. Since increasing the strength of a steel sheet can degrade its formability, the development of a material with both high strength and high formability has been anticipated. A high-strength steel sheet has been proposed that utilizes strain-induced transformation of retained austenite, exhibiting high strength and excellent ductility. This steel sheet has a microstructure that includes retained austenite and can be easily formed into a predetermined shape due to the retained austenite during forming. The steel sheet then achieves high post-forming strength as a result of the transformation of retained austenite into martensite. For example, Patent Literature 1 proposes a high-strength steel sheet with a tensile strength of 1000 MPa or more, a total elongation (EL) of 30% or more, and markedly high ductility. This steel sheet is manufactured using a strain-induced transformation of retained austenite. The steel sheet is produced by causing a steel sheet containing C, Si, and Mn as fundamental constituents to form austenite, subsequently quenching the steel sheet in the bainite transformation temperature range, and performing isothermal holding, i.e., an austempering treatment. The retained austenite forms as a result of the concentration of C in the austenite due to the austempering treatment, and the formation of a large amount of retained austenite requires the addition of a large amount of C exceeding 0.3%.However, an increase in carbon concentration in steel degrades spot weldability, and in particular, a carbon concentration above 0.3% significantly degrades spot weldability. Therefore, it has been difficult to bring this technique into active use for automotive steel sheets. Furthermore, the ability to be flanged and bent is not considered in the patent literature because the primary objective is to improve the ductility of the high-strength steel sheet. IVIA / a / ¿U¿¿ / UU4DO l In Patent Literature 2, a high-Mn steel is used, and a suitable strength-ductility balance is achieved by performing heat treatment in the ferrite-austenite dual-phase region. However, no study is conducted in Patent Literature 2 on improving ductility by concentrating Mn in untransformed austenite. Therefore, there is room for improvement in workability. In Patent Literature 3, a medium-Mn steel is used, and the total elongation is increased by performing a heat treatment in the ferrite-austenite dual-phase region, concentrating Mn in untransformed austenite and thus forming stable retained austenite. However, the achievement of compatibility between elongation, flangability, and bendability is not considered. Furthermore, no study has been conducted on improving flangability and bendability by controlling the distribution of not only Mn but also C in the second phase composed of retained austenite and martensite. The Mn concentration is considered insufficient to achieve compatibility between elongation, flangability, and bendability because the heat treatment time in the manufacturing method described in Patent Literature 3 is short, and the Mn diffusion rate is low. Furthermore, in Patent Literature 4, a medium-Mn steel is used, and uniform elongation and flangability are improved by heat-treating a hot-rolled steel sheet in the ferrite-austenite dual-phase region for an extended period. This facilitates the concentration of Mn in untransformed austenite, thus forming retained austenite grains with a high aspect ratio. However, in Patent Literature 4, the improvement of ductility and flangability of a high-strength steel sheet is studied solely through Mn concentration; no study is conducted on the compatibility of improving flangability, bendability, and elongation by controlling the C and Mn distribution in the second phase composed of retained austenite and martensite. List of Appointments Patent Literature PTL 1: Publication of unexamined Japanese patent application no. S61-157625 PTL 2: Publication of unexamined Japanese patent application no. H1-259120 PTL 3: Publication of unexamined Japanese patent application no. 2003-138345 PTL 4: Japanese Patent No. 6123966 Brief Description of the Invention Technical Problem The present invention was made in light of the circumstances described above. An object of the present invention is to provide a high-strength steel sheet having a tensile strength (TS) of 980 MPa or more and excellent formability, and a method for manufacturing the high-strength steel sheet. Note that the term formability used herein refers to ductility, drawability, and bendability. Solution to the Problem To achieve the above object, the inventors of the present invention carried out extensive studies in terms of the chemical composition of the steel sheet and a method for manufacturing the steel sheet to manufacture a high-strength steel sheet with excellent formability and, consequently, found the following facts. Specifically, the inventors of the present invention found it important to heat a steel slab including Mn at a content of 2.50% by mass or more and 8.00% by mass or less and, optionally, other alloying elements, such as, at a suitable content, hot roll the steel slab to a rolling delivery temperature of 750°C or more and 1000°C or less, then perform winding at 300°C or more and 750°C or less, optionally perform holding at a temperature equal to or lower than the transformation temperature Aci for more than 1800 s, then perform cold rolling, subsequently perform holding at a temperature equal to or higher than the transformation temperature Acs of -50°C for 20 s more and 1800 s less, then perform cooling to a cooling stop temperature equal to or lower than the martensitic transformation start temperature,Then, reheat to a reheating temperature of 120 °C or higher and 450 °C or lower, subsequently hold at the reheating temperature for 20 seconds or more and 600 seconds or less, and then cool to room temperature in such a way that film-like austenite grains, to which carbon is concentrated and which serve as nuclei for fine retained austenite, form grains with a high aspect ratio in the subsequent annealing step. After cooling, hold at a temperature equal to or higher than the transformation temperature Aci and equal to or lower than the transformation temperature Aci of +150 °C for 20 seconds or more and 600 seconds or less, and then cool. It was found that the method described above allows the production of a high-strength steel sheet with excellent formability, the steel sheet having a steel microstructure that includes,by area, ferrite: 30% or more and 80% or less, martensite: 5% or more and 35% or less, and retained austenite: 8% or more, wherein the ratio of the area fraction of retained austenite grains, grains having an aspect ratio of 2.0 or more and a minor axis length of 1 mm or less, divided by the total area fraction of retained austenite is 0.3 or more, wherein the ratio of the average Mn content (mass %) in retained austenite divided by the average Mn content (mass %) in ferrite is 1.5 or more, and the product of the ratio of the average Mn content (mass %) in retained austenite divided by the average Mn content (mass %) in ferrite and the average aspect ratio of retained austenite is 3.0 or more, the ratio of the average C content (% by mass) in retained austenite divided by the average C content (% by mass) in ferrite is 3.0 or more,and where the ratio of the average C content (% by mass) in the retained austenite divided by the average Mn content (% by mass) in the retained austenite is 0.05 or more. The present invention was made based on the results described above. The brief description of the present invention is as follows. [1] A high-strength steel sheet comprising a chemical composition containing, by mass, C: 0.030% or more and 0.250% or less, Si: 0.01% or more and 3.00% or less, Mn: 2.50% or more and 8.00% or less, P: 0.001% or more and 0.100% or less, S: 0.0001% or more and 0.0200% or less, N: 0.0005% or more and 0.0100% or less, and Al: 0.001% or more and 2.000% or less, the remainder being Fe and incidental impurities; a steel microstructure comprising, by area, ferrite: 30% or more and 80% or less, martensite: 5% or more and 35% or less, and retained austenite: 8% or more, wherein a ratio of a grain area fraction of the retained austenite, the grains having an aspect ratio of 2.0 or more and a minor axis length of 1 μπΊ or less, divided by a total area fraction of the retained austenite is 0.3 or more, wherein a ratio of an average Mn content (mass %) in the retained austenite divided by an average Mn content (mass %) in the ferrite is 1.5 or more, and a product of the ratio of the average Mn content (% by mass) in the retained austenite divided by the average Mn content (% by mass) in the ferrite and an average aspect ratio of the retained austenite is 3.0 or more, wherein a ratio of the average C content (% by mass) in the retained austenite divided by an average C content (% by mass) in the ferrite is 3.0 or more, and wherein a ratio of the average C content (% by mass) in the retained austenite divided by the average Mn content (% by mass) in the retained austenite is 0.05 or more. [2] The high-strength sheet steel is described in [1], wherein the chemical composition further contains at least one element selected from, by mass, 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, Ca: 0.0050% or less, Mg: 0.0050% or less, Zr: 0.0050% or less, and REM: 0.0050% or less. [3] The high-strength steel sheet described in [1] or [2], the high-strength steel sheet which further includes a galvanized layer disposed on a surface of the high-strength steel sheet. [4] The high-strength steel sheet described in [3], wherein the galvanized layer is a galvanized annealed layer. [5] A method for manufacturing a high-strength steel plate, the method comprising heating a steel slab having the chemical composition described in [1] or [2], hot rolling the steel slab to a finish roll delivery temperature of 750°C or more and 1000°C or less, then coiling at 300°C or more and 750°C or less, then cold rolling, then holding at a temperature equal to or greater than an Acs transformation temperature of -50°C for 20 seconds or more and 1800 seconds or less, then cooling to a quenching stop temperature equal to or less than a martensitic transformation start temperature, then reheating to a reheating temperature of 120°C or more and 450°C or less, then holding at the reheating temperature for 2 seconds or more and 600 seconds or less, then cooling to room temperature,subsequently perform maintenance at a temperature equal to or higher than the Acr transformation temperature equal to or lower than the Aci transformation temperature of +150 °C for 20 seconds plus or minus 600 seconds, and then perform cooling. [6] The method for manufacturing a high-strength steel sheet described in [5], the method which further includes, after winding and before cold rolling, holding at a temperature equal to or lower than the Acr transformation temperature for more than 1800 s. [7] The method for manufacturing a high-strength steel sheet described in [5] or [6] also includes performing a galvanizing treatment. [8] The method for manufacturing a high-strength steel sheet described in [7], the method which further includes, after galvanizing treatment, performing an alloying treatment at 450 °C or more and 600 °C or less. Advantageous Effects of the Invention According to the present invention, a high-strength steel sheet can be produced that has a tensile strength (TS) of 980 MPa or more and is excellent in terms of formability, specifically, not only ductility but also flexibility and stretchability. The application of a high-strength steel sheet produced by the manufacturing method according to the present invention, for example, to structural components of automobiles reduces the weight of car bodies and thus improves fuel efficiency. Therefore, the use of the high-strength steel sheet is highly valuable from an industrial perspective. Description of the Modalities The present invention is specifically described below. Hereinafter, % used to describe the content of the elements means % by mass, unless otherwise specified. (1) The reasons why the content of the steel components is limited to the above ranges in the present invention are described below. C: 0.030% or more and 0.250% or less Carbon (C) is a necessary element for forming low-temperature transformation phases, such as martensite, and therefore increasing strength. C is also an effective element for improving the stability of retained austenite and the ductility of steel. If the C content is less than 0.030%, it is difficult to achieve the expected area fraction of martensite, and the expected strength may not be reached. Furthermore, it is difficult to achieve a sufficient area fraction of retained austenite, and adequate ductility may not be achieved. If C is added to steel in an excessive amount, such that the content If the carbon content (C) exceeds 0.250%, the hard martensite area fraction increases to an excessive level. In such a case, during a hole expansion test, the number of microvoids formed at the martensite grain boundaries may increase, and crack propagation may proceed unfavorably. Therefore, the drawability may be degraded. Furthermore, the weld zone and the heat-affected zone may harden significantly, and consequently, the mechanical properties of the weld zone may be degraded. Therefore, spot weldability, arc weldability, and similar properties may be degraded. From the points of view described above, the carbon content is limited to 0.030% or more and 0.250% or less. The carbon content is preferably 0.080% or more and 0.200% or less. Yes: 0.01% or more and 3.00% or less It is effective in achieving adequate ductility, as it improves the ferrite's resistance to deformation. If the Si content is less than 0.01%, the advantageous effects of Si addition may be minimal. Therefore, the lower limit is set at 0.01%. Adding Si to steel in an excessive amount, such that the Si content exceeds 3.00%, can cause embrittlement of the steel, which degrades ductility and bendability, and the formation of red scale, which degrades surface quality. Furthermore, coating quality may be degraded. Consequently, the Si content is limited to 0.01% or more and 3.00% or less. The preferred Si content is 0.20% or more and 2.00% or less, and preferably 0.20% or more and less than 0.70%. Mn: 2.50% or more and 8.00% or less Manganese (Mn) is an extremely important additional element in the present invention. Mn stabilizes retained austenite and is effective in achieving adequate ductility. Mn also increases the strength of steel by strengthening solid solutions. Furthermore, Mn is effective in forming stable retained austenite in which it concentrates, thus achieving the intended ratio of the average Mn content (mass %) in retained austenite to the average Mn content (mass %) in ferrite, and consequently, adequate ductility. The actions described above are confirmed when the Mn content in the steel is 2.50% or higher. However, if excessive amounts of Mn are added to the steel, such that the Mn content exceeds 8.00%, a non-uniform microstructure may form due to Mn segregation, and consequently, the drawability may be degraded.Furthermore, the expected ratio of the average carbon content (mass %) in retained austenite to the average manganese content (mass %) in retained austenite may not be achieved. This makes it difficult to achieve adequate ductility and bendability. Additionally, the ease of conversion treatment and the quality of the coating may be compromised. From the perspectives described above, the manganese content is limited to 2.50% or more and 8.00% or less. The manganese content is preferably 3.10% or more and 6.00% or less, and more preferably 3.20% or more and 4.20% or less. iviA / a / zuzz / uu^oo i P: 0.001% or more and 0.100% or less Phosphorus (P) is an element that strengthens the solid solution and can be added to steel according to the desired strength. P also facilitates ferrite transformation and is therefore effective in forming a multiphase microstructure. To achieve the advantageous effects described above, the P content must be limited to 0.001% or more. If the P content exceeds 0.100%, weldability may be compromised. Furthermore, if zinc coating alloying is performed, the alloying rate may be reduced, and the quality of the zinc coating may be degraded. Consequently, the P content is limited to 0.001% or more and 0.100% or less, and preferably 0.005% or more and 0.050% or less. S: 0.0001% or more and 0.0200% or less Sulfur (S) segregates at grain boundaries, causing brittleness in steel during hot working, and is present as sulfides, which degrade local deformability. Consequently, it is necessary to limit the S content to 0.0200% or less. The S content is preferably 0.0100% or less, and more preferably 0.0050% or less. However, due to production technology limitations, it is necessary to limit the S content to 0.0001% or more. Therefore, the S content is limited to 0.0001% or more and 0.0200% or less. The content is preferably 0.0001% or more and 0.0100% or less, and more preferably 0.0001% or more and 0.0050% or less. N: 0.0005% or more and 0.0100% or less Nitrogen (N) is an element that degrades the aging resistance of steel. This degradation becomes particularly significant if the N content exceeds 0.0100%. Although minimizing the N content is preferable, it is necessary to limit it to 0.0005% or more due to production technology limitations. Consequently, the N content is limited to 0.0005% or more and 0.0100% or less. The preferred N content is 0.0010% or more and 0.0070% or less. To: 0.001% or more and 2.000% or less Aluminum (Al) is an element that broadens the ferrite-austenite dual-phase region and is effective in reducing the dependence of mechanical properties on the annealing temperature, thus improving the stability of those properties. The lower limit is set at 0.001% because the advantageous effects of adding Al are minimal if the Al content is below 0.001%. Al also acts as a deoxidizing agent and is effective in improving the cleanliness index of steel. Therefore, it is preferable to add Al to steel during the deoxidation step. However, adding excessive amounts of Al to steel, such that the Al content exceeds 2,000%, increases the risk of cracking in steel slabs during continuous casting and reduces manufacturability. Based on the above considerations, the Al content is limited to 0.001% or more and 2,000% or less. The Al content is preferably 0.200% or more and 1.200% or less. In addition to the constituents described above, the chemical composition may contain at least one element selected from, by mass, 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, Ca: 0.0050% or less, Mg: 0.0050% or less, Zr: 0.0050% or less, and REM (short for rare earth metals): 0.0050% or less. Ti: 0.200% or less Titanium (Ti) is effective for precipitation strengthening of steel. Ti increases the strength of ferrite, thus reducing the hardness difference between ferrite and the second hard phase (martensite or retained austenite) and consequently allowing for a more suitable drawability. When Ti is added to steel, the Ti content is preferably 0.005% or more, and ideally 0.010% or more. However, if the Ti content exceeds 0.200%, the area fraction of hard martensite increases to an excessive level. In such a case, during a hole expansion test, the number of microvoids formed at the grain boundaries of the martensite may increase, and crack propagation may proceed unfavorably. Therefore, the drawability may be degraded. Consequently, when Ti is added to steel, the Ti content is limited to 0.005%.200% or less. The Ti content is preferably 0.100% or less. Nb: 0.200% or less, V: 0.500% or less, and W: 0.500% or less Nitrogen (Nb), vanadium (V), and tungsten (W) are effective in precipitation strengthening of steel. Furthermore, similar to the advantageous effects of adding titanium (Ti), Nibrois (Nb), V, and W, they increase ferrite strength, reducing the hardness difference between ferrite and the second hard phase (martensite or retained austenite) and consequently allowing for greater flangability through proper drawing. When Nibrois, V, and W are added to steel, the contents are preferably 0.005% or more, and 0.010% or more, respectively. However, if the Nb content exceeds 0.200% or the V or W content exceeds 0.500%, the hard martensite area fraction increases excessively.In such a case, during a hole expansion test, the number of microvoids formed at the martensite grain boundaries may increase, and crack propagation may proceed unfavorably. Therefore, the drawability may be degraded. Consequently, if Nb is added to the steel, the Nb content shall be limited to 0.200% or less. The Nb content is preferably 0.100% or less. If V or W is added to the steel, the V or W content is limited to 0.500% or less. The V or W content is preferably 0.300% or less. B: 0.0050% or less Boron inhibits the formation and growth of ferrite at austenite grain boundaries. Boron increases the strength of ferrite, thus reducing the hardness difference between ferrite and the second hard phase (martensite or retained austenite) and consequently enabling greater flangability. IVIA / a / ¿U¿¿ / UU4D01 most suitable stretching. In the event that B is added to the steel, the B content is preferably 0.0003% or more. The B content is more preferably 0.0005% or more. However, if the B content exceeds 0.0050%, formability may be degraded. Consequently, in the event that B is added to the steel, the B content shall be limited to 0.0050% or less. The B content is preferably 0.0030% or less. Ni: 1,000% or less Ni is an element that stabilizes retained austenite and is effective in achieving more suitable ductility. Ni also increases the strength of steel by strengthening solid solutions. When Ni is added to steel, the Ni content is preferably 0.005% or more. However, if the Ni content exceeds 1.000%, the area fraction of hard martensite increases excessively. In such a case, during a hole expansion test, the number of microvoids formed at the martensite grain boundaries may increase, and crack propagation may proceed unfavorably. Therefore, the drawability may be degraded. Consequently, when Ni is added to steel, the Ni content is limited to 1.000% or less. Cr: 1.000% or less and Mo: 1.000% or less Cr and Mo can be added to steel as needed because they improve the balance between strength and ductility. When Cr and Mo are added to steel, the Cr and Mo contents are preferably 0.005% or more and 0.005% or more, respectively. However, if Cr and Mo are added to steel in excessive amounts, such that the Cr content exceeds 1.000% and the Mo content exceeds 1.000%, the hard martensite area fraction increases to an excessive level. In such a case, during a hole expansion test, the number of microvoids formed at the martensite grain boundaries may increase, and crack propagation may proceed unfavorably. Therefore, the drawability may be degraded. Consequently, if the above elements are added to steel, the content of the elements is limited to Cr: 1,000% or less and Mo: 1,000% or less. Cu: 1,000% or less Copper (Cu) can be added to steel as needed because it is an effective steel-strengthening element. If copper is added to steel, the copper content is preferably 0.005% or higher. However, if copper is added to steel in an excessive amount, such that the copper content exceeds 1.000%, the area fraction of hard martensite increases to an excessive level. In such a case, during a hole expansion test, the number of microvoids formed at the martensite grain boundaries may increase, and crack propagation may proceed unfavorably. Therefore, the drawability may be degraded. Consequently, if copper is added to steel, the copper content is limited to 1.000% or lower. Sn: 0.200% or less and Sb: 0.200% or less iviA / a / zuzz / uu^oo i Tin (Sn) and antimony (Sb) are added to steel as needed to inhibit decarburization of a region of the steel sheet's surface layer that is several tens of micrometers thick, resulting from nitriding or oxidation of the steel sheet surface. Sn and Sb are effective in inhibiting this nitriding or oxidation, preventing a reduction in the martensite area fraction on the steel sheet surface and thus achieving a degree of strength and stability in mechanical properties. When Sn and Sb are added to steel, the Sn and Sb content is preferably 0.002% or higher. However, if either element is added to steel in excessive amounts, such that the element content exceeds 0.200%, toughness may be degraded. Consequently, when Sn and Sb are added to steel, the Sn and Sb content is limited to 0.200% or lower. Ta: 0.100% or less Similar to Ti and Nb, Ta increases strength by forming carbide and carbonitride alloys. Furthermore, Ta is considered to partially dissolve in Nb carbide or carbonitride to form a compound precipitate, such as (Nb,Ta)(C,N), which significantly reduces precipitate oiling and stabilizes the strength increase through precipitation strengthening. Therefore, Ta is preferably added to steel. When Ta is added to steel, its content is preferably 0.001% or more. However, adding excessive amounts of Ta to steel can saturate the precipitation stabilization effect and increase alloying costs. Consequently, when Ta is added to steel, its content is limited to 0.100% or less. Ca: 0.0050% or less, Mg: 0.0050% or less, Zr: 0.0050% or less, and REM: 0.0050% or less Calcium (Ca), magnesium (Mg), zinc (Zr), and mineral spirits (ME) are effective elements for increasing the sphericity of sulfides and reducing their adverse effects on drawability. If these elements are added to steel, their content should preferably be 0.0005% or higher. However, if any of these elements is added to steel in excessive amounts, such that the element content exceeds 0.0050%, the amount of inclusions and similar defects may increase, and consequently, surface and internal defects may occur. Therefore, if Ca, Mg, Zr, and ME are added to steel, their contents should each be limited to 0.0050% or lower. The component other than the components described above includes Fe and incidental impurities. (2) The microstructure of steel is described below. Ferrite Area Fraction: 30% or more and 80% or less It is necessary to limit the ferrite area fraction to 30% or more to achieve sufficient ductility. It is also necessary to limit the soft ferrite area fraction to 80% or less to achieve a tensile strength of 980 MPa or more. Note that the term ferrite used herein refers to polygonal ferrite, granular ferrite, and acicular ferrite, i.e., ferrite materials that are relatively soft and ductile. The ferrite area fraction is preferably 40% or more and 75% or less. Martensite Area Fraction: 5% or more and 35% or less It is necessary to limit the martensite area fraction to 5% or more to achieve a tensile strength (TS) of 980 MPa or more. It is also necessary to limit the martensite area fraction to 35% or less to achieve adequate ductility and bendability. The martensite area fraction is preferably 5% or more and 30% or less. The term martensite used herein refers to quenched martensite. The area fractions of ferrite and martensite can be determined by grinding a cross-section of the thickness (L-section) of the steel sheet that is parallel to the rolling direction, etching the L-section in 3% nital, and then observing the 1 / 4 thickness position (i.e., the position of 1 / 4 of the steel sheet thickness below the surface in the depth direction) of the L-section with a SEM (scanning electron microscope) at 2000x magnification for 10 fields of view, calculating the area fractions of the microstructure components (ferrite and martensite) with Image-Pro produced by Media Cybemetics, Inc. for each of the 10 fields of view based on the resulting microstructure images, and taking the averages thereof.In the microstructure images above, ferrite appears as a gray microstructure component (ground microstructure) and martensite appears as a white microstructure. Fraction of Retained Austenite Area: 8% or more It is necessary to limit the retained austenite area fraction to 8% or more to achieve sufficient ductility. The retained austenite area fraction is preferably 12% or more and 25% or less. The retained austenite area fraction was determined by grinding the steel sheet to a position 0.1 mm below the 1 / 4 thickness position, further grinding the steel sheet 0.1 mm by chemical polishing, and then measuring the integrated intensity ratios of the diffraction peaks in the {200}, {220}, and {311} planes of fcc iron and in the {200}, {211}, and {220} planes of bcc iron with an X-ray diffraction apparatus using CoKa radiation, and taking the averages of the nine integrated intensity ratios. The ratio of the area fraction of retained austenite grains having an aspect ratio of 2.0 or more and a minor axis length of 1 mm or less divided by the total area fraction of retained austenite is 0.3 or more Limiting the ratio of the area fraction of retained austenite grains having an aspect ratio of 2.0 or more and a minor axis length of 1 pm or less divided by the total area fraction of retained austenite at 0.3 or more is an important condition constituting the present invention. Retained austenite grains having an aspect ratio of 2.0 or more and a minor axis length of 1 pm or less reduce the likelihood of void formation during punching performed prior to the flange forming step and thus improve the draw-flangeability. To achieve adequate draw-flangeability, the area fraction of retained austenite grains having an aspect ratio of 2.0 or more and a minor axis length of 1 pm or less is limited.The aspect ratio should be 0 or higher, and the minor axis length should be 1 pm or less. Furthermore, the retained austenite area fraction should be sufficiently high to achieve high ductility. The aspect ratio should preferably be 0.5 or higher. The upper limit for the aspect ratio should preferably be 15.0 or less. The lower limit for the minor axis length should preferably be 0.05 pm or more, which is the detection limit in EBSD. Martensite and retained austenite were identified using a phase map obtained by EBSD (electron backscattered diffraction). The aspect ratio of a retained austenite grain was calculated by drawing an ellipse circumscribing the retained austenite grain using Photoshop 13 and dividing the length of the ellipse's major axis by the length of its minor axis. Ratio between the average Mn content (% by mass) in the retained austenite and the average Mn content (% by mass) in the ferrite: 1.5 or more Limiting the ratio of the average Mn content (mass %) in retained austenite to the average Mn content (mass %) in ferrite to 1.5 or greater is an extremely important condition that constitutes the present invention. To achieve adequate ductility, the fraction of stable retained austenite area in which Mn is concentrated must be high. The above ratio is preferably 2.0 or greater. Although the upper limit for the above ratio is not restricted because the higher the average Mn content in the retained austenite, the greater the ductility, the above ratio is preferably 10.0 or less because the increase in ductility saturates if the above ratio exceeds 10.0. The product of the ratio between the average Mn content (% by mass) in the retained austenite divided by the average Mn content (% by mass) in the ferrite and the average aspect ratio of the retained austenite grains is 3.0 or more Limiting the product of the ratio of the average Mn content (mass %) in retained austenite divided by the average Mn content (mass %) in ferrite and the average aspect ratio of the retained austenite grains to 3.0 or higher is extremely important. To achieve adequate ductility, the area fraction of stable retained austenite grains with a high aspect ratio where Mn is concentrated must be high. The above index is preferably 4.0 or higher. The upper limit for the above index is 20.0 or lower. Ratio of the C content (% by mass) in the retained austenite divided by the average C content (% by mass) in the ferrite: 3.0 or more Limiting the ratio of the carbon content (mass %) in the retained austenite to the average carbon content (mass %) in the ferrite to 3.0 or more is an extremely important condition that constitutes the present invention. When the carbon content in the retained austenite is increased, the retained austenite remains even in a high-stress region during bending deformation, and consequently, bendability is improved. Therefore, to achieve adequate ductility and bendability, the fraction of stable retained austenite area in which Mn is concentrated must be high. The above ratio is preferably 5.0 or more. The upper limit for the above ratio is preferably 10.0 or less. The ratio of the average C content (% by mass) in the retained austenite divided by the average Mn content (% by mass) in the retained austenite is 0.05 or more Limiting the ratio of carbon content (mass percent) in retained austenite to manganese content (mass percent) in retained austenite to 0.05 or greater is an extremely important condition that constitutes the present invention. To achieve adequate bendability as a result of carbon concentration in retained austenite, the fraction of stable retained austenite area in which not only manganese but also carbon is concentrated must be high. When both carbon and manganese are concentrated in retained austenite, stable retained austenite can be present in a high-stress region, and consequently, suitable ductility and bendability compatibility can be achieved. The above ratio is preferably 0.06 or greater. The upper limit for the above ratio is preferably 0.10 or less. The carbon (C) and manganese (Mn) content in retained austenite and ferrite is determined using a three-dimensional atomic probe (3DAP) with a sample taken from the 1 / 4-thickness position. First, a portion of the steel sheet containing retained austenite and ferrite is sampled, and then an acicular sample is formed using a focused ion beam. A voltage is applied to the acicular sample with the 3DAP, and the released carbon (C) and manganese (Mn) ions are analyzed. The manganese (Mn) content can be determined as a percentage by dividing the number of atoms of C and Mn measured by the total number of atoms measured for each retained austenite and ferrite grain. This measurement is performed for each of the 30 retained austenite grains and 30 ferrite grains randomly selected from the field of view, and the averages of the C and Mn contents determined by the quantitative analysis are calculated.The C and Mn content (% by mass) in retained austenite and ferrite can be obtained by converting the C and Mn content (% by atom) on a mass basis. The advantageous effects of the present invention are not impaired even if the microstructure of the steel according to the present invention contains tempered martensite, bainite, pearlite or carbides, such as cementite, in addition to ferrite, martensite and retained austenite, when the area fraction of tempered martensite, bainite, pearlite or carbides is 10% or less. The steel sheet may include a galvanized layer applied to the surface. The galvanized layer may be an annealed galvanized layer formed by alloying the galvanized layer. (3) The manufacturing conditions are described below. Temperature at which the steel slab is heated The temperature to which the slab is heated is preferably, among others, 1100 °C or higher and 1300 °C or lower. Since the precipitates present when the steel slab is heated are present as coarse precipitates in the final steel sheet and do not affect strength, the Ti and Nb-based precipitates formed during casting can be redissolved. It is also preferable to limit the temperature to which the steel slab is heated to 1100 °C or higher to reduce air bubbles, segregation, and similar features present in the surface layer of the slab, further reduce cracks and irregularities present on the surface of the steel sheet, and thus further flatten the surface of the steel sheet. The temperature to which the steel plate is heated is preferably 1300 °C or lower to reduce the loss of scale caused by increased oxidation.The heating temperature of the slab above is preferably 1150 °C or more and 1250 °C or less. Steel slabs are preferably manufactured by a continuous casting process to avoid macrosegregation. Alternatively, steel slabs can be manufactured by ingot casting, thin slab casting, or similar methods. In addition to the conventional method where, after production, the steel slab is cooled to ambient temperature and then reheated, energy-saving processes such as hot-charged rolling and direct hot rolling can be used without problems. In these processes, the steel slab is not cooled to ambient temperature but is instead charged into a heating furnace while its temperature is high, or it is rolled immediately after a short period of thermal insulation. The slab is formed into a sheet bar by rough rolling under ordinary conditions.If the heating temperature is relatively low, the foil bar is preferably heated with a bar heater or similar before determining the rolling to avoid the occurrence of problems during hot rolling. Hot rolling delivery temperature 750 °C or more and 1000 °C or less The heated steel slab is hot-rolled to form a hot-rolled steel sheet through rough rolling and finish rolling. If the delivery temperature of the previous finish rolling exceeds 1000 °C, the amount of oxides (flakes) formed increases rapidly, the roughness of the interface between the base iron and the oxides increases accordingly, and the surface quality may degrade after pickling and cold rolling. Furthermore, if the hot-rolled flakes remain partially unmoved after pickling, they negatively affect ductility and drawability. In addition, the grain size may increase to an excessive level, and the surface roughness of a pressed article may increase during working.If the delivery temperature of the previous finishing roll is below 750 °C, the rolling force increases, the rolling load increases accordingly, and the proportion of reduction performed while the austenite is in a non-recrystallized state increases. In such a case, an abnormal texture develops. Consequently, the in-plane anisotropy of the final product increases significantly. This degrades the uniformity of the material quality (stability of mechanical properties). Furthermore, the aspect ratio of the retained austenite grains is reduced. IVIA / a / ¿U¿¿ / UU4DO l As a consequence, ductility and drawability may be degraded. Therefore, it is necessary to limit the delivery temperature for hot rolling to 750 °C or higher and 1000 °C or lower. The delivery temperature for finishing rolling is preferably 800 °C or higher and 950 °C or lower.Temperature at which winding is performed after hot rolling: 300 °C or higher and 750 °C or lower. If the temperature at which winding is performed after hot rolling exceeds 750 °C, the grain size of the ferrite included in the microstructure of the hot-rolled steel sheet increases, the aspect ratio of the retained austenite grains included in the final annealed steel sheet decreases, it becomes difficult to achieve the intended ratio of the area fraction of retained austenite grains with an aspect ratio equal to or greater than 2.0 and a minor axis length equal to or less than 1 mm divided by the total area fraction of retained austenite, and the drawability is consequently degraded. If the temperature at which winding is performed after hot rolling is lower than 300 °C, the strength of the hot-rolled steel sheet increases.In such cases, the rolling load required for cold rolling may increase, and defects in the shape of the steel sheet may occur. This reduces productivity. Consequently, it is necessary to limit the temperature at which winding is performed after hot rolling to 300 °C or higher and 750 °C or lower. The winding temperature prior to hot rolling is preferably 400 °C or higher and 650 °C or lower. Finish rolling can be performed continuously by joining the as-rolled steel sheets together during hot rolling. The as-rolled steel sheets can be temporarily coiled. To reduce the rolling force required for hot rolling, some or all of the finish rolling can be performed using a lubricant. Lubrication rolling is also preferable to increase uniformity in the shape of the steel sheet and the uniformity of material quality. When lubrication rolling is performed, the coefficient of friction is preferably 0.10 or higher and 0.25 or lower. Hot-rolled steel sheet produced as described above can be optionally pickled. Pickling is preferable because it removes oxides from the surface of the steel sheet and further improves the ease of conversion treatment and the quality of the coating. Pickling can be carried out in one or more steps. Maintaining a temperature equal to or lower than the Aci transformation temperature for more than 1800 seconds It is preferable to perform the holding process at or below the transformation temperature Aci for more than 1800 seconds because it softens the steel sheet that will be cold-rolled in the subsequent step. If the holding process is performed at a temperature above the transformation temperature Aci, the Mn concentrates in the austenite, hard martensite and retained austenite form after cooling, and the steel sheet may not soften. In such a case, retained austenite may form at the grain boundaries in the subsequent annealing step, increasing the amount of retained austenite grains with a low aspect ratio. This makes it difficult to achieve the intended area fraction of retained austenite grains with an aspect ratio of 2.0 or more and a shaft length of less than 1 pm or less divided by the fraction of the total area of ​​the retained austenite, and consequently, may degrade the drawability. Even at a temperature equal to or lower than the Aci transformation temperature, if the holding time is 1800 or less, it is difficult to eliminate the stress remaining after hot rolling, and the steel sheet may not soften. Heat treatment can be performed using any annealing method, such as continuous or batch annealing. After heat treatment, cooling is carried out at room temperature. There are no limitations on the cooling method or rate. Any cooling method can be used, such as furnace cooling or natural cooling for batch annealing, or gas jet cooling, mist cooling, or water cooling for continuous annealing. If pickling is required, a conventional method can be used. Cold Rolling The resulting steel sheet is cold-rolled. The cold-rolling reduction ratio is preferably between 15% and 80%. Performing cold rolling at the above reduction ratio allows for the formation of a sufficiently recrystallized microstructure and further improves ductility. Maintenance at a temperature equal to or higher than the Acs transformation temperature of -50°C for 20 seconds plus and 1800 seconds minus If holding is performed at a temperature below the Acs transformation temperature of 50 °C, the Mn concentrates in the austenite, martensitic transformation does not occur during cooling, and consequently, retained austenite grain nuclei with a high aspect ratio may not form. In such a case, in the subsequent annealing step, retained austenite may form at the grain boundaries in a disadvantageous manner. This increases the number of retained austenite grains with a low aspect ratio and makes it impossible to form the intended microstructure. Even at a temperature equal to or above the Acs transformation temperature of -50 °C, if the holding time is less than 20 s, recrystallization does not occur to a sufficient degree, and the intended microstructure does not form. Consequently, ductility may be degraded.If the previous holding is carried out for more than 1800 s, the amount of hard martensite increases, and ductility and bendability may degrade accordingly. Cooling temperature to stop cooling equal to or lower than the temperature at which the martensitic transformation begins. If the cooling stop temperature is higher than the martensitic transformation start temperature, when the amount of martensite to be transformed is small, all the untransformed austenite may transform into martensite during the final cooling, and no retained austenite grain nuclei with a high aspect ratio are formed. In such a case, in the subsequent annealing step, the retained austenite may form at the grain boundaries in a disadvantageous manner. This increases the amount of retained austenite grains with a low aspect ratio and makes it impossible to form the intended microstructure. The cooling stop temperature is preferably equal to or higher than the martensitic transformation start temperature of -250 °C and equal to or lower than the martensitic transformation start temperature of -50 °C. Reheating to a reheat temperature of 120 °C or more and 450 °C or less and holding at the reheat temperature for 2 seconds more and 600 seconds less If the reheating temperature is below 120 °C, carbon does not concentrate in the retained austenite formed in the subsequent annealing step, and consequently, the intended microstructure does not form. If the reheating temperature exceeds 450 °C, the nuclei of retained austenite grains with a high aspect ratio decompose, and the amount of retained austenite grains with a low aspect ratio increases accordingly. This makes it impossible to form the intended microstructure. Similarly, if the holding time is less than 2 seconds, the nuclei of retained austenite grains with a high aspect ratio cannot form, and the intended microstructure cannot form.If the retention time exceeds 600 s, the nuclei of retained austenite grains with a high aspect ratio decompose, and the number of retained austenite grains with a low aspect ratio increases accordingly. This makes it impossible to form the intended microstructure. Cooling is carried out at room temperature after reheating and holding. The cooling method is not limited. For example, natural cooling, gas cooling, or mist cooling at 200 °C or lower followed by water cooling is preferable. Pickling can be carried out as needed. If pickling is performed, common collection methods can be used. Maintenance at a temperature equal to or higher than the Aci transformation temperature and equal to or lower than the Aci transformation temperature of +150°C for 20 seconds plus and 600 seconds minusPerforming maintenance at a temperature equal to or higher than the transformation temperature Aci and equal to or lower than the transformation temperature Aci of +150 °C for 20 seconds plus or minus 600 seconds is an extremely important condition that constitutes the present invention. If maintenance is performed at a temperature lower than the transformation temperature Aci for less than 20 seconds, the carbides formed during heating remain undissolved, making it difficult to form sufficient area fractions of martensite and retained austenite, and consequently reducing strength. If the aforementioned maintenance temperature exceeds the transformation temperature Aci of +150 °C, the area fraction of martensite increases, and consequently, the intended microstructure does not form. Furthermore, Mn and C do not concentrate in the austenite to sufficient levels, and consequently, ductility and bendability are degraded.Furthermore, a sufficient fraction of retained austenite area does not form, and ductility degrades accordingly. The holding temperature mentioned is preferably equal to or lower than the Aci transformation temperature of +100 °C. If the holding time exceeds 600 s, the martensite area fraction increases, and consequently, the intended microstructure does not form. Additionally, Mn and C do not concentrate in the austenite to sufficient levels, and consequently, ductility and bendability degrade. Furthermore, the fraction of retained austenite area sufficient to achieve the desired ductility does not form. Galvanizing Treatment In the case of hot-dip galvanizing, the steel sheet that has undergone the previous annealing treatment is immersed in a galvanizing bath at a temperature of 440 °C or higher and 500 °C or lower for hot-dip galvanizing. The coating weight is then adjusted by gas cleaning or a similar method. For hot-dip galvanizing, a galvanizing bath with an aluminum content of 0.08% or higher and 0.30% or lower is preferably used. Electrogalvanizing or a similar treatment may be used instead of hot-dip galvanizing. If zinc coating alloying is performed after galvanizing (C), the coating alloying treatment is carried out at 450°C or higher and 600°C or lower. If the alloying treatment is performed at higher than 600°C, the untransformed austenite may transform into pearlite, and the intended fraction of retained austenite area may not be achieved. Consequently, ductility may be degraded. Therefore, if zinc coating alloying is performed, it is preferable to carry it out at 450°C or higher and 600°C or lower. Although the other manufacturing method requirements are not limited, it is preferable to perform the aforementioned annealing treatment using a continuous annealing facility from a productivity standpoint. It is also preferable to perform the entire set of treatments—annealing, hot-dip galvanizing, and zinc coating—using a continuous galvanizing line (CGL), which is a hot-dip galvanizing line. The above high-strength steel sheet and the high-strength galvanized steel sheet, which includes the high-strength steel sheet and a galvanized coating on the surface, can be subjected to pass rolling for purposes such as shape correction and surface roughness adjustment. The rolling reduction ratio in pass rolling is preferably 0.1% or more and 2.0% or less. If the above rolling reduction ratio is If the reduction ratio (IVIA / a / ¿U¿¿ / UU4D01) is less than 0.1%, the advantageous effects are small and it is difficult to control the reduction ratio. Therefore, this is considered the lower limit for a suitable rolling reduction ratio. If the rolling reduction ratio exceeds 2.0%, productivity may be significantly degraded. Therefore, this is considered the upper limit for a suitable rolling reduction ratio. Pass rolling can be performed online or offline. Pass rolling with a specified reduction ratio can be performed in one or more passes. In addition, various coating treatments can be used, such as resin coating and grease coating. EXAMPLES The molten steels with the chemical compositions described in Table 1, the remainder being Fe and incidental impurities, were prepared using a converter and formed into slabs by a continuous casting process. The slabs were reheated to 1250 °C and then formed into high-strength (CR) cold-rolled steel sheets with a thickness of 1.0 to 1.8 mm under the conditions described in Tables 2 and 3. A galvanizing treatment was then applied to produce hot-dip galvanized (Gl) steel sheets. The hot-dip galvanized steel sheets were subsequently alloyed to produce hot-dip galvanized (GA) steel sheets. The hot-dip galvanizing bath used to produce the hot-dip galvanized (Gl) steel sheets was a zinc bath containing 0.19% Al by mass.The hot-dip galvanizing bath used to produce the hot-dip galvanized (HDG) steel sheets was a zinc bath containing 0.14% Al by mass. The bath temperature was 465°C. The coating weight was 45 g / m² per side (both sides were coated). In the production of the HDG steel sheets, the Fe concentration in the coated layer was adjusted to 9% by mass or more and to 12% by mass or less. The cross-sectional steel microstructure, tensile properties, and drawability of each of the aforementioned steel sheets were determined. The results are listed in Tables 4, 5, and 6. Table 1 iviA / a / zuzz / uu^oo i Steel Type Chemical Composition (% by mass) Ms Temperature (=C) Transformation Temperature Aci (=C) Transformation Temperature Ac: (=C) Observed c Si Mn PSN Al Ti Nb VWB Ni Cr Mo Cu Sn Sb Ta Ca Mg Zr RE MA 0.1 65 0.55 3.53 0.021 0.0022 0.0034 0.030 0.049 -352 656 772 Invention Steel B 0.1 84 0.76 3.22 0.026 0.0025 0.0041 0.047 0.041 -358 666 786 Invention Steel c 0.1 77 1.80 3.59 0.0 19 0.00 19 0.00 21 0.0 34 0.0 33 345 667 818 Invention Steel D 0.2 41 0.98 3.30 0.0 29 0.00 29 0.00 25 0.0 58 - 335 666 767 Invention Steel E 0.0 48 1 00 4.10 0.0 31 0.00 24 0.00 26 0.0 31 - 370 646 794 Invention Steel F 0.1 80 2.90 3.97 0.0 27 0.00 19 0.00 26 0.0 31 0.0 44 - 329 669 859 Invention Steel AH 3.1 60 43 1 OK 20 0 2Ό 21 000 26 33 9Í - - - - - - - - - 51 - - - - 365 662 794 Invention steel Al 3.1 33 is} Ϊ3 3 03 9 003 30 30 - - - - - - - - - - - 06 - - - 361 656 766 Invention steel AJ 3.2 01 40 c'c 3Í 300 20 000 26 29 - 06 361 663 753 Invention steel AK Ϊ2 22 2 23 3.03 27 000 40 32 π - 03 32S 646 7?o Invention steel AL ϊθ g'é 24 3 03 25 0 03 40 40' - 3 00 32 31S 647 749 Steel of invention AV 3.1 95 22 '5 22 0 20 24 0 03 36 35 - - - - - - - - - - - - - OCO 24 - 331 655 769 Steel of invention AN 40 03 0 Cl 26 3_00 0 00 23 40 08 - 0 00 32 347 663 733 Steel of invention AC 79 05 Í2 2Í 3 03 28 0 03 36 4Ϊ - o_oo ¿ < 279 579 68C Steel of invention Part sut'ayaca Outside of anb to from to present rvercon,- Centén-co to nrve of imp j'ezas ncidentales IVIA / a / ¿U¿¿ / UU4D01 The starting temperature of the martensitic transformation, i.e., the Ms temperature, and the transformation temperatures Aci and Aca were calculated using the following formulas. Onset temperature of martensitic transformation: Temperature Ms (°C) = 550 - 350 x (% of C) - 40 χ (% of Mn) -10 χ (% of Cu) - 17 χ (% of Ni) - 20 χ (% of Cr) -10 χ (% of Mo) - 5 % of V χ (% of W) + 30 χ (% of Al) Transformation temperature Aci (°C) = 751 - 16 χ (% of C) + 11 χ (% of Si) - 28 χ (% of Mn) - 5.5 χ (% of Cu) -16 χ (% of Ni) + 13 χ (% of Cr) + 3.4 χ (% of Mo) - 28 χ (% of Mn) - Transformation temperature Ac3(°C) = 910 - 203 χ (% of C) + 45 χ (% of Si) - 30 χ (% of Mn) - 20 χ (% of Cu) -15 χ (% of Ni) + 11 χ (% of Cr) + 32 χ (% of Cr) + 32 χ (% of Mo deV) + 400 χ (% ofT) + 200 χ (% of Al) where: (% of C), (% of Si), (% of Mn), (% of Ni), (% of Cu), (% of Cr), (% of Mo), (% of V), (% of Ti), (% of V), (% of W) and (% of Al) each represent the contents (% by mass) of the element. Table 2 No. Tp 0 ce ace ro Delivery temperature ce a laminaoón final rCi Bob naco temperature rCi Heat treatment of hot-rolled steel sheet Reduction coefficient of 9 ladina cor en *ríc Annealing treatment of cold-rolled steel sheet Annealing treatment of cold-rolled steel sheet Alloy temperature reí Tip O* Heat treater temperature rCi Heat treatment time (51 Heat treatment temperature rCi Heat treatment time 0 (5.1 Cooling stop temperature rCi Reheating temperature !·^ Holding time of the reheating temperature Reheating temperature thermal rCi Thermo treatment time 0 'SI 1 M seo 520 540 130C0 A* · V soo •5C 125 300 340 700 •5C CR 2 M S80 530 515 234C0 52 9 S50 'SC 180 325 250 700 •se Gl 3 M seo 460 580 2S-SC0 56.3 S20 •2C 150 250 120 685 •5C 500 G4 4 M eco 450 550 234C0 54 .7 S05 ’5C 120 200 250 710 30C 550 GA A M 910 520 500 144C0 52 5 610 ‘5C 200 350 150 680 •2C Gl A M 920 5C0 600 1S0C0 53 3 900 10 240 280 140 690 •3C 520 G4 M 870 540 625 18OCO 57 * 780 2400 80 120 250 710 •2C 510 8 n SCO 550 570 360C0 529 750 2ÜC 410 420 180 730 •5C Gl 6 M S20 550 56.3 SCO 25C 300 500 215 745 80 480 1 c A S50 4¿0 600 144C0 64 .7 SOO •2C 50 1CO 300 665 60C 480 G7 1 1 A seo 520 500 seco 5’ * S05 50 120 180 900 700 4SC Gl 1 2 A 375 380 535 9CO0 A* · V S20 26C 240 260 1. 690 •se CR 1 3 A 950 460 462 S35 25C 180 275 500 680 •4C 540' 1 4 B seo 5CC> 580 2-6C0 54 3 S40 S2C 320 420 540 710 •5C Gl 1 5 C 910 520' 550 2'6C0 529 S50 •5C 250 300 180 705 •5C 550 G4 1 A A ¿co 540 550 1S0C0 47. ‘ S80 'SC 140 215 210 740 25C Gl 1 A seo S50 590 360C0 56 5 S50 24C 180 220 120 670 •se 510 1 8 A S70 020 750 2*6C0 53 3 S50 •se 110 200 6C 700 •5C Gl 1 6 A S50 560 430 1GSC0 53 3 SCO 3OC 120 195 370 520 2OC 500 G4. c A 880 540! 540' 18OC0 63 4 790' 300 175 330 530 S5C1 450 490' GA 1 A SCO 550 520' 7200 6' * 780' '50 150' 180 170 680' 10' 525 GA 2 2 A 870 580 550 234C0 84 7 750' -se 110' 250 260 710' 2000 540 GA 3 D 905 580 550' 2SSC0 53 3 880' 1203 140 260 180 700' •20 CE 4 ES¿0 550' 570' 180C0 58 8 SCO1 360 180' 310 240 680' 500 Gl |ru to | F 875 600' 530' 18000 57 * 850' -5C 180' 295 540 750' 90' CR AG SSO 810 490' 23400' 57. 330' -4C 100' 250 120 700' 240 550' . GA 2 H 375 5CC' 525 324C0 53 3 340' '20 150' 320 270 680' 90' 510' GA 8 I 910 580' 530 234C0 50.0 775 -5C 150' 340 570 715. 3-3C 530' GA ns J 880 5CC' 510' 2SSC0 529 780' •se 130' 300 30 690' •30 CR 3 CK 870 420' 520' 324C0 48.6 790' 90' 60 200 220 650' •20 Gl 3 1 L 950 580' 550' 234C0 462 780' •oc 170' 240 150 725 720 Gl 3 2 V 880 610 580' 36000 62.5 720' 50 130' 250. 150 695 •20 550' GA 3 3 N SSO 6C0 530' 28-300' 52.5 SCO' '20 130' 210 180 700' •se CR Underlined portion: Outside the scope of the present invention * CR: Cold-rolled steel sheet (without coating), Gl: Hot-dip galvanized steel sheet (without zinc coating), GA: Hot-dip galvanized steel sheet Table 3 κ c P 0 of steel Final rolling delivery temperature Winding temperature 0 'heat treatment of hot rolled steel sheet Cold rolling reduction coefficient Annealing treatment of cold rolled steel sheet Annealing treatment of cold rolled steel sheet *η'o Alloy temperature Tp c Heat treatment temperature Tier-p c of heat treatment 0 (5) Heat treatment temperature Tier-p c of heat treatment 0 is.i Cooling parachute temperature Reheating temperature Holding time of reheating temperature Heat treatment temperature Time c of heat treatment 0 (5) 34 3 830 520 520 10800 58 8 6'0 '30 250 355 120 765 •80 500 GA 35 P 86-0 430 52 C 630 40 150 20C 400 790 360 520 GA 36 Q 870 560 600 9CO0 56 3 630 300 2Ό 325 80' 735 '80 CR 37 R 855 550 58 8 620 '40 2-0 30C 90' 630 50' 5*5 GA 38 850 850 550 7200 62 5 870 625 130 20C 19C 7O •25540 GA 36 830 600 540 1O8CO 62 5 830 720' 240 325 1SC 700 '80' CR 4C T 850 430 430 10800 64 7 8'0 60 50 isc 190 7Ό 300' Gl 41 u eoo 540 520 350CO 57 1 830 '00' 240 35C too 725 340 Gl 4¿ *4 860 600 550 238CO 5C C 8'0 330' 250 40C 500 700 •80' CR “2 w 6'0 500 56 3 890 '20' 200 33C 140 695 ‘80' 520 GA 44 X 600 550 5'0 353CO 46 2 840 '50' 130 32C 225 730 600 π tí) GA 45 Ύ 870 550 570 14400 52 9 850 40' 55 125 170 890 80' CR 46 z 505 330 47 1 825 300' 130 30C 310 700 250' 5Ό GA 4? AA 890 6'0 530 238CO 55 6 820 1200 30*0 405 260 695 60 520 GA 48 AS 6'0 540 530 13000 56 3 840 40' 140 isc 270 725 200' Gl 4g AC 870 740 520 234CO 58 8 800 60 120 15C 160 750 250' Gl 50 AD 835 6'0 590 2'SCO 53 3 600 240 130 35C 100 720 •80' 530 GA 51 AE 830 500 520 234CO 64 7 8'0 ‘20' 140 1SC 210 700 60 Gl 52 AF 500 500 570 9CO0 62 5 830 '50' 1O0 21C 140 7Ό 250' 50*0 GA 53 AG 6'0 530 5'0 238CO 56 3 840 '50' 200 32C 210 740 30 5Ό GA 54 AH 855 530 53 8 820 60' 150 175 1SC 740 340 520 GA 55 A 500 560 520 324CO 56 3600 •40' 95 isc 190 7-5 •80' Gl 56 AJ 903 553 543 10803 56 3 903 90 103 200 125 883 '70' 530 GA 57 853 553 543 14403 56 3 803 80 183 21C 150 7'5 240' CR 58 AL 883 523 5'3 10803 64 7 825 '50 173 190 150 603 '50' 480 GA 59 AV 843 503 46 7 853 '80 103 300 240 703 •80' CR 60 AN 853 403 56-3 2'603 50 C 843 240 173 305 180 683 210' 543 GA 91 AO 883 503 5'5 9CO3 57 1 835 320 123 2?5 300 653 •50' 505 GA IVIA / a / ¿U¿¿ / UU4D01 Underlined part: Outside the scope of the present invention * GR: Cold-rolled steel sheet (without coating), Gl: Hot-dip galvanized steel sheet (without zinc coating alloy), GA: Hot-dip galvanized steel sheet The tensile test was performed in accordance with JIS Z 2241 (2011) using a JIS test specimen n.s5 prepared by taking a sample from each of the steel sheets such that the tensile direction was perpendicular to the rolling direction of the steel sheet. In the tensile test, TS (tensile strength) and EL (total elongation) were measured. Regarding the mechanical properties, a "Good" rating was given in the following cases. When TS was 980 MPa or more and less than 1080 MPa, EL > 20% When TS was 1080 MPa or more and less than 1180 MPa, EL > 16% The drawability was measured in accordance with JIS Z 2256 (2010). Specifically, each steel sheet was cut into a 100 mm x 100 mm piece. A 10 mm diameter hole was formed in the piece with a clearance of 12% ± 1%. Subsequently, while the piece was held by a die with an inner diameter of 75 mm and a clamping force of 9 tons, a tapered punch with a vertex angle of 60 degrees was inserted into the hole. The diameter of the hole was then measured at the critical point where cracking occurred. A critical hole expansion ratio λ (%) was calculated using the following formula. The drawability was assessed based on this critical hole expansion ratio. Critical hole expansion ratio λ (%) = {(Df - Do) / Do} × 100 where Dt represents the diameter (mm) of the hole in which cracking occurred, and Do represents the initial diameter (mm) of the hole. In the present invention, a Good rating was given in the following cases, depending on the TS. When TS was 980 MPa or more and less than 1080 MPa, λ> 15% When TS was 1080 MPa or more and less than 1180 MPa, λ > 12% The bendability test was performed in accordance with the V-block method described in JIS Z2248 (1996), using a bendability test specimen 30 mm wide and 100 mm long taken from each of the annealed steel sheets, with the rolling direction parallel to the bending axis (bending direction). The test was performed at a stroke speed of 100 mm / syn = 3 for each bending radius. The presence of cracks on the outside of the bent portion was determined using a stereomicroscope. The minimum bending ratio at which no cracking occurred was considered the critical bending ratio R. In the present invention, a "Good" rating was given for the bendability of the steel sheet if the critical bending ratio at 90° VR / t < 2.5 was met (t: thickness of the steel sheet). The high-strength steel sheets prepared in exemplary inventions had a tensile strength (TS) of 980 MPa or more. In exemplary inventions, high-strength steel sheets with excellent formability were prepared. In contrast, in comparative examples, at least one of the properties (TS, EL, λ) and the tensile bending capacity was poor. iviA / a / zuzz / uu^oo i Table 4 '4:· T 20 thousand Thickness ;ΤΓΤΊ; rsozion de ares de F “raocion de ares de M í'.; F-acccn de ares ce PA i%! Fracccn de ares ce PA con -elation de specio iocsl c supe-ion a 2.0 y longitud ce eje nrenc·- >qual 0 infeno- a 1 m Fraction de a-ea coca de Mído 1%. in msssi Ccntenib 0 p-orrede ce 01 r eF i:-: in mass: Oo-ienido average of Mn in RA come- do average of Mn M 1 4 co c --TTv- 2U.o 2 351 5.ob —ΤΤΓ- 3 10 M ó 53! 79 2.33 2 43 M 4 545 20.5 Ή : 5J2 5.45 2.78 2 32 4 M 4 49.7 23.6 17.6 0 550 5.28 2.74 2 25 EM 2 53.0 170 91 9'. 5.55 2.16 3 04 r¡ M • 4 23.5 15.3 21.3 0 4'3 5 27 3.12 1 59 M 2 ​​41 2 29.4 17.3 w 3-^ 5.75 2.30 2 4' 54' 249. 1?5 5 12 2.46 2 45 s M 4 51 2 25 2 15.9 C ¿Co 5 84 2.51 20' 10 M 4 15.2 20.3 7 497 8.7 > 2.63 3 33 11 M • 45 5 17.3. 201 5.34 2.88 2 23 12 M 4 5TM C 13.1 19.5 2227 5 10 2.31 2 17 13 M ' 4 55 c 13.3 20.2 0 5'2 7.22 1.71 4 14 E2 4 •3. 20.2 22.0 0 593 5 05 1.25 4 04 15 517 13.5 24.3 0 454 5 12 1.31 3 38 15 - c 55.7 23.3 18.2 0 ¿D5 5 53 2.43 2 27 17 • ó 50 5 19.9 15.4 0 253 5.01 0.88 5 8'. 18 M • 4 43.3 27 3 19.9 2 321 6.45 2.93 2 15 10 M ' 4 73.5 3d 7 1 2 353 4 55 2.56 1 7d 20 M 2 53.1 35.1 — 4 25 2.03 1 43 21 M ' 4 75.5 3 o — w oζ < 5 92 3.C4 1 95 22 M 2 36 5 35 9 55 3 7 5 4 15 2.97 1 40' 23 J • 4 60 4 13.9 18.5 2 54C 5.39 2.71 1 99 24 E ' 4 53.2 14.5 23.3 0 4'4 4 55 2.30 1 53 25 F 2 τ 79 7 19.1 2 393 4 14 2.10 1 97 26 G 2 46 4 23.9 15.4 2 35C 6.77 2Λ4 2 7^ H • 4 56 4 21.0 12.8 2 5'7 *0.01 3.10 3 23 2δ ' 4 53.2 24 6 15.3 2 7* 1 739 2.30 2 73 20 J 53.4 13.7 25.3 2 35C 4 67 η π*? 2K 30 k 8 49.4 22.1 2 0 751 * 121 4.45 2 47 31 L 4 53 5 20.4 15.3 0 3-3^ 6.24 2.35 2 17 32 M 2 51.0 21.1 23.3 0 5'4 d .25 2.45 33 N 2 R¿ 5 10.3 17.1 2 391 5 99 02 2 05 34 0 • 4 KP 5 20. 0 10.6 2 3* 1 7 30 3.20 2 44 35 P 2 50.3 13.5 25.0 0 421 7.01 1.01 4 15 35 Q 4 49.9 23 3 14.3 r. ~ k 9 5 95 1.06 5 6* 37 5 4 70.3 4δ 7 1 w 0 V 5 52 2.45 2 25 38 e 2 S» A 143 25.1 2 577 7 15 2.65 2 73 30 C 2 53.2 14.5 25.9 0 556 6.88 2.0δ 2 3 40 2 49.6 15.2 28.4 0 3'6 9 95 d.12 1 63 ¿1 J ¿- 55 5 23.6 6 8 2 420 3.73 Π 7 1 44 42 ν • 4 5^ 5 36 5 δ 1 0 452 4 95 η 1 9 2 33 43 W ' 4 55 3 17.1 26.7 0 5'5 5.32 3.10 2 04 44 19.C 10.6 2 3*9 4 73 2.99 1 57 45 52 6 19.9 24.4 2 3*0 5.24 2.41 2 17 45 8 53.2 15.3 20.0 2 3'6 5 12 3.C2 1 73 ΑΑ 51.1 22. C 23.7 2 399 5.25 2.56 2 05 4δ ΑΒ • 4 40 5 13.2 20.3 2' 479 007 2.93 3 40 40 AC • 4 50.1 216 2 421 *0.43 4.21 2 48 50 • 4 43 2 31.0 20.3 0· 4'0 7.29 3.03 24* 51 ΑΕ 2 54' 1?7 22.3 2 393 6 14 2.10 2 92. 52 AF 2 5+0 17.5 22.3 2 &53 5.71 1.93 2 95 5S AG • 4 rn 2 21.5 21.3 0 422 4 93 3.02 1 55 54 Mr* 4 49.3 33 C 13.6 '582 5 05 2.93 1 73 55 AJ ' 4 43 5 30 C 15.8 2 3+4 4.35 2.74 1 59 ob ' 4 511 20.1 20.4 2' 437 5.50 3.30 1 72' Oí AK ' 4 49.5 20 5 20.1 3 323 5.43 3.28 1 95 55 AL 4 50.1 20.1 20.5 2 321 5 83 2.52 2 83 50 AM w* 53.1 19.9 24.3 0 375 5.41 3.03 1 dü ' 4 Kn 7 20.2 20.8 w' 0 □ 4.7.· 1.71 2 79 of AO 4 513 20.5 21.0 5'4 '2.25 4.50 2 53 IVIA / a / ¿U¿¿ / UU4D0i Part below: Outside the scope of this invention F: Ferrite, M: Martensite, RA: Retained Austenite Table 5 V* i Ί c:n¡5 :>.· RA Ü-cilftidt: de Mn •jf i RA L'j-'cr ti:: ;rct·ι“·: u r<· Mr cr el ' Rt-.r ::r •redi;i de dsix't.iu :v RA C.j-er d:: jrcrirdc de C en RA । ‘ ·.· : C:: [.-o· ιΐυ'ΐκΐυ v.lr:· :!v »_.·· » r t-is.í υ::·Ί7-··Ί::··.; j.*»j*i d ¿ <·ν RA Ü::n!?ri |.- / rr !i·:. C «η Γ μ·α·ι <*: u d;t - 'T AR i:-.*i>.—rx» μυτν::ι::· Λν \'ί c- ΛΡ •Jr tiq. rjLür.i il ur i Ó’ 42 IT. ¿i ! -·- .1 Uf; :: 4.! ’l· B ~ P - * .. .· 144 4 1 .1 LU 1¿t: Jt '1.· B « 4 írd 1* i‘: .1 U!; 1: > ' '1 / B P ή * j j 1 · -· .1 1 ’ 4 Lh· ·>' '1.· B ►· Π 4 44 1.! 4?: :. 44 .i L»r '1 / Ü ►' A| ·> J 'j < 4 :. !4 .1 l.U 1: .ir. '1.· B P «j * .. J 1.. M ’ :. .14 .1 ll·· '1 / B ►· i i. .! <’ .1 Uí: • L: .ir ’l· B ►' 4 Ί i .1 Uf: :! d4 .>' '1.· B P 7 1.: i r.<¡ .1 Li-· !: Li’ ¿i '1.· B ► 7 1 * 1 .4 .. -·.· .1 lU r z- '1,· B P 5 u i 4r. .1 U7 '1 / B ►· 1.! ·’ - X ’.í^ :. 44 .i i.: 4 4.: .ir. '1 / L‘ A, 14 ·· ·: > 1 7 <Jl. >. - .· .1 i · 4 Ll·: ’l· B P ñ lü j 1 * '4 ’ 4o .1 £ i- ’! / B P A 1C i .. - .1 Ut: 1,· B r 2 t > 1' <r¿ ’¿ 40 0 Ü2 f ü.: ü’ M 0 ** 7 1 2 i % 4 R I. ’.l 12 4 l.lj ’.l’: ’l.· t< ► ” 1 : - '· •J Ik! 4 i.».: ’J 1 ’f,· D ► X -· 1 t. .· > . .! :· '.l 1.! 4 *2«^ '.i;. 7 / B - 4 tO J 1.! '.l 14 1 7 '.ih '1.· B • - 7 l7 ji '.l 14 T., 'VB * -· > 2o 1.: / :. '· -* '.l U ; ·« x 2 '.ir ”h· B x-<j 4 1 i.: Ή ’.l U4 i ? ’.ii: 'Μ B · — ,.c • - i.: ui. jm > .i i.: ·?.: '.i'·. 'MD · “ i '20 i.: 77 · Cj ¿1 Üf: 1.: í:.<ii: 'M D P * - 4 7‘.l C Ü1 '.I Li4 14 .··: ’.ll: ’l· B r — - ..«» 4!: > -7 '.i ir. 2 1 ' '.ir. 'l· B · 7 -- 5 2'-J 1 · '.I··. •_l l«2 p.m. 4j „ 1 'l· B 7 4 to 1 · 51: •.1 U2 7. x'· '.i': '1.· B 7 '.i : > 1.1 L 4M '.l 114 14 2': '.ii: '1.· B 5 4 o 1:! 4 ¡ '.l 1 · 4 U4 '.l': '1.· B 7 jJ 1 ·- I7 .17 ;. ri '.l U 5 1.: '.i1. 'L· B · 7 o 4 5 if 72 '.l U2 ri '.l': 'L· B 7· 7 ixo 1 Ά „ 44 'J> '.ir. D ► 7 4 ;..3 x ' '44 · ·— •. Γ: > 2 '.l': 'Μ B ► 7 * -1 '.l - '.l L· ¿2 '.L. .. '.l.: 'ML' *- 7 • - 1.! o * '.l 1'4 1:! Ht: .. 'J.: 'l· B 7 -·- 4 44 1.: 4': 1. 1 '.l L1.! if 'j .. ’J . ’M B * 7 4.: O 4 i.: >4 .. - .· ’.l U ’ 7. '· · >.i4 ’M B · 7 4' · i •j l.f” wX ’.i u5 :! 4.: 1 ’M B ► 7 4. ? 44 14 oí: ’.l 1 · 4 ’.r Γ.. '1.· B ► 7 4j - x :> 4..' 4i.i ’.i i? 5 21».: ’Jh '1.· B ► 7 44 • ·- >. - < ’.i ir. 7 M· 1:. '1.· B * 7 4f 4 2'J < ¿ 5u ’.i ir. 2 '> i 1:. ’L· B 7· 7 4í. • - o r»:; · r ’.i ir. 2 ó7 1:. _h· B 7 4 ' ’.l 12 oh • r ; ’.l L·:· r. o ! i:. D * 7 4:! x'.: ot: ;. 4r. ’.l L·:· Γ -r ’.H: D * 7 4:' - ..*) r ¡:. ’J US r o.: ’.ll: ’M Ü 7 > 44 1.! ’.l·. ’.l Li4 i ’.ih 'Μ B r 7 > 2 i 1f > - -7 ’.l UJ 1f. _¡ J ’.i!: ’l· B ► 7. MA / a / XUXX / UU^OO1 n1 1' 7>: i. j ir. T Dr p >4 5 hy M 7 ' L 41 J 'Jtr - ...t! '.l·. 'L· Β A - uJ uf: .1 Mr 4 UJ '.11 *1.' D- ► >4 ~S ·· :> ir. - J .1 U!: t* O.. .: r„ '1,- B-' r >4 4 7>-J ί 1.! 1. .! jr .! 11! oír ML! ► <4 >.) 1.: 5': L 4 1 .1 Mr - .· un '1,- B '4· 4 · '4 · '4. .1 U7 u? ').· Β ' >4 ? .. - .1 M:! b.! ilt: π· ► >4 Ul 1J L 4 1 .1 U: f 1.! SL ; > ·, .1 LL · '.Γ: '1,- B -· ' 4. iviA / a / zuzz / uu^oo i Underlined part: Outside the scope of the present invention F: Ferrite, TM: Tempered Martensite, RA: Retained Austenite BF: Bainitic ferrite, P: Pearlite, 0: Carbides (e.g., cementite) Table 6 No. (MPa) EL % λ fmmj Rt Observations • 995 22 5 22 3 4 2 4 Example of irvencicr 2 1021 24 9 21 3 6 2 3 Example of irvencicr o 1038 26 C 22 3 2 2 3 Example of irvencicr 4 1097 21 2 19 2 4 20 Example of irvencicr 5 1025 25 9 12 26 22 Comparative example 6 1054 14 C 27 O. Z 2 3 Comparative example 1051 12 1 18 26 44 Comparative example O 1'75 142 1C 2 4 1.5 Comparative example 9 T7C 12 5 9 3 4 2 4 Comparative example 1C 102C 23 8 21 5 2 3.3 Comparative example 11 1054 18 3 14 3 4 2 4 Comparative example 12 985 19 1 14 3 2 2 3 Comparative example 13 1031 23 9 24 1 8 1.3 Irvencicr example 14 1052 22 4 23 3 0 2 1 Irvencicr example 15 1016 27 1 25 2 8 1.S Irvencicr example 16 1'32 14 2 11 2 8 1.6 Comparative example 17 1013 24 8 13 2 0 2.0 Comparative example 18 984 25 8 11 3 4 2 4 Comparative example ie 045 26 0 28 1 0 0.7 Comparative example 2C 1'54 13 δ 14 3 0 3.0 Comparative example 21 9C2 25 4 3C 1 2 oa Comparative example 22 1096 13 7 26 3 8 3.2 Comparative example 23 1045 2C 0 17 20 1.a Invention example 24 08· 22 4 25 3 4 2 4 Invention example 25 1'6C 22 3 16 1 4 1 2 Example of irvenciór 20 1078 23 5 16 2 8 2.3 Example of irvenciór 27 1032 2C 6 20 0 8 0 0 Example of irvenciór 28 1030 24 2 18 3 0 2.1 Example of irvenciór 26 089 20 4 18 2 8 18 Example of irvenciór 3C 1005 27 8 23 3 4 1.0 Example of irvenciór 31 1084 21 C 2C 2 5 18 Example of irvenciór 32 1O5C 23 6 22 2 4 2.0 Example of irvenciór 33 1023 22 2 21 3 0' 2 5 Example of irvenciór 34 066 21 5 2C 3 2 2.3 Example of irvenciór 35 1023 23 C 2C 2 4 2.0 Example of irvenciór 30 1030 21 5 18 2 8 2.0 Example of irvenciór 37 842 15 3 27 2 4 17 Comparative example 38 060 15 C 27 3 4 2.8 Comparative example 36 085 15 1 21 3 5 2.a Comparative example 4C 1032 14 1 14 3 0 3.0 Comparative example 41 1025 13 δ 16 2 4 2.0 Comparative example ¿2 1022 24 3 13 20 1.a Comparative example 43 1013 20 1 22 2 0' 1.4 Example of irvenciór 44 1 3C 23 2 23 2 0 1.0 Example of irvenciór 45 069 22 8 2C 2 8 1.8 Example of irvenciór 40 089 21 5 22 2 8 1.0 Example of irvenciór ¿7 1074 16 8 25 2 4 1.5 Example of irvenciór 48 082 3C 3 2C 2 8 2.0 Example of irvenciór 46 1013 21 C 16 3 2 2.3 Example of irvenciór 5C 1032 2C 3 18 - - 1.0 Example of irvenciór. 51 1O3C 21 4 23 2 8 2.3 Irvenciór plc exercise 52 965 24 5 24 1 8 1.5 Irvenciór exercise 53 1'04 24 3 22 1 4 1.0 Irvenciór example 54 1087 21 2 2C 3 0 2.5 Irvenciór example 55 1034 22 4 16 26 1.9 Irvenciór exercise 56 967 26 6 22 3 0 2.1 Irvenciór example 57 984 23 1 23 26 1.9 Irvenciór example 58 1026 22 5 24 2 4 2.0 Irvenciór exercise irvenciór 56 966 26 7 22 3 0 1.9 Example of irvenciór 6C 963 21 1 24 2 4 1.7 Example of irvenciór 61 985 22 C 23 3 0 2 5 Example of irvenciór IVIA / a / ¿U¿¿ / UU4D01 Underlined portion: Outside the scope of the present invention Industrial Applicability According to the present invention, a high-strength steel sheet can be produced that has a tensile strength (TS) of 980 MPa or more and excellent formability. The application of the high-strength steel sheet according to the present invention, for example, to structural components of automobiles reduces the weight of car bodies and thus improves fuel efficiency. Therefore, the use of the high-strength steel sheet according to the present invention is highly valuable from an industrial perspective.

Claims

1. A high-strength steel sheet comprising: a chemical composition containing, by mass, C: 0.030% or more and 0.250% or less, Si: 0.01% or more and 3.00% or less, Mn: 2.50% or more and 8.00% or less, P: 0.001% or more and 0.100% or less, S: 0.0001% or more and 0.0200% or less, N: 0.0005% or more and 0.0100% or less, and Al: 0.001% or more and 2.000% or less, the remainder being Fe and incidental impurities; a steel microstructure, including, by area, ferrite: 30% or more and 80% or less, martensite: 5% or more and 35% or less, and retained austenite: 8% or more, wherein a ratio of a grain area fraction of retained austenite, the grains having an aspect ratio of 2.0 or more and a minor axis length of 1 gm or less, divided by a total area fraction of retained austenite is 0.3 or more, wherein a ratio of the average Mn content (% by mass) in the retained austenite divided by an average Mn content (% by mass) in the ferrite is 1.5 or greater, and a product of the ratio of the average Mn content (% by mass) in the retained austenite divided by the average Mn content (% by mass) in the ferrite and an average aspect ratio of the retained austenite is 3.0 or greater, wherein a ratio of an average C content (% by mass) in the retained austenite divided by an average C content (% by mass) in the ferrite is 3.0 or more, and wherein a ratio of the average C content (% by mass) in the retained austenite divided by the average Mn content (% by mass) in the retained austenite is 0.05 or more.

2. The high-strength steel sheet according to claim 1, wherein the chemical composition contains at least one element selected, by mass, 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, Ca: 0.0050% or less, Mg: 0.0050% or less, Zr: 0.0050% or less, and REM: 0.0050% or less.

3. The high-strength steel sheet according to claim 1 or 2, the high-strength steel sheet further comprises a galvanized layer disposed on a surface of the high-strength steel sheet.

4. The high-strength steel sheet according to claim 3, wherein the galvanized layer is a galvanized annealed layer.

5. A method for manufacturing a high-strength steel sheet, the method comprising heating a steel slab having the chemical composition according to claim 1 or 2, hot rolling the steel slab to a finish rolling delivery temperature of 750°C or more and 1000°C or less, then coiling at 300°C or more and 750°C or less, then cold rolling, then holding at a temperature equal to or greater than a transformation temperature (Acs) of -50°C for 20 seconds or more and 1800 seconds or less, then cooling to a holding temperature equal to or less than a martensitic transformation start temperature, then reheating to a holding temperature of 120°C or more and 450°C or less, then holding at the holding temperature for 2 seconds or more and 600 seconds or less, then cooling to room temperature.Subsequently, perform maintenance at a temperature equal to or greater than the transformation temperature Aci equal to or less than the transformation temperature Aci of +150 °C for 20 seconds plus or minus 600 seconds, and then perform cooling.

6. The method for manufacturing a high-strength steel sheet according to claim 5, the method further comprising, after winding and before cold rolling, holding the sheet at a temperature equal to or lower than the transformation temperature Aci for more than 1800 s.

7. The method for manufacturing a high-strength steel sheet in accordance with claim 5 or 6, the method further comprising performing a galvanizing treatment.

8. The method for manufacturing a high-strength steel sheet according to claim 7, the method further comprising, after galvanizing treatment, performing an alloying treatment at 450 °C or more and 600 °C or less.