HIGH-STRENGTH STEEL SHEET AND METHOD FOR MANUFACTURING IT

Abstract

The objective is to provide a high-strength steel sheet having a yield strength greater than 0.70, a tensile strength of 980 MPa or more, and excellent formability, as well as 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, tempered martensite: 3.0% 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 1 µm 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 the retained austenite divided by the average Mn content (mass %) in the ferrite is 1.5 or more, and the product of the ratio of the average Mn content (mass %) in the retained austenite divided by the average Mn content (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 (mass %) in the retained austenite divided by the average C content (mass %) in the ferrite is 3.0 or more, and wherein the ratio of the average C content (mass %) in the retained austenite divided by the average Mn content (mass %) in the retained austenite is 0.05 or more, and wherein the diffusible hydrogen content in the steel is 0.3 ppm in mass or less.

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 high yield ratio (YR) of more than 0.70 and a tensile strength (TS) of 980 MPa or more, 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 allow for 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. If a steel sheet with a tensile strength (TS) of 980 MPa or higher is used as a vehicle frame member, it is required to have high formability and a high yield strength (YR) to protect vehicle occupants. 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. 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, in Patent Literature 2, no study is conducted 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 Mn concentration is considered insufficient to achieve the intended compatibility of elongation, yield strength, drawability, and bendability, since the amount of time during which the heat treatment is performed 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 drawability of a high-strength steel sheet is studied solely through Mn concentration; no study is conducted on improving the creep coefficient and bendability 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 high yield ratio (YR) of more than 0.70, 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 objective, the inventors of the present invention carried out extensive studies regarding the chemical composition of the steel sheet and a method of manufacturing the steel sheet and, consequently, ascertained the following facts. Specifically, the inventors of the present invention found it important to limit the Mn content to 2.50% by mass or more and 8.00% by mass or less, adjust the contents of the other alloying elements, such as Ti, to be within suitable ranges, optionally hold at a temperature equal to or lower than the transformation temperature Aci for more than 1800 s after hot rolling, then perform cold rolling, then hold at a temperature equal to or higher than the transformation temperature Acs of -50 °C for 20 s more and 1800 s less, and then cool to a cooling stop temperature equal to or lower than the martensitic transformation start temperature, so that film-like austenite grains are formed, serving as the core of fine, retained austenite grains with a high aspect ratio, in the subsequent reheating stage.Once cooling has stopped, reheating is performed to a reheating temperature equal to or higher than the transformation temperature of Aci and equal to or lower than the transformation temperature Am of +150 °C. Subsequently, the reheating temperature is held for 20 minutes or more and 1800 minutes or less, and then cooling is performed. Additionally, galvanizing may be carried out as required. Furthermore, alloying may be performed at 450 °C or higher and 600 °C or lower. Subsequently, cooling is performed to 100 °C or lower, followed by holding at 100 °C or higher and 400 °C or lower for 10 minutes or more, and then cooling is performed. It was found that the method described above allows the production of a high-strength steel sheet with excellent formability, a steel sheet with a steel microstructure that includes, by area, ferrite: 30% or more and 80% or less, martensite: 3%.0% or more and 35% or less, and retained austenite: 8% or more, wherein the ratio of the grain area fraction of the 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 the retained austenite is 0.3 or more, wherein the ratio of the average Mn content (mass %) in the retained austenite divided by the average Mn content (mass %) in the ferrite is 1.5 or more, and the product of the ratio of the average Mn content (mass %) in the retained austenite divided by the average Mn content (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 (mass %) in the retained austenite divided by the average content of C (% by mass) in 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 Mη content (% by mass) in the retained austenite is 0.05 or more, and wherein the diffusible hydrogen content in the steel is 0.3 ppm by mass or less. 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 pm 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, and wherein a diffusible hydrogen content in the steel is 0.3 ppm by mass or less. [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 sheet comprising a steel microstructure comprising, by area, ferrite 30% or more and 80% or less, quenched martensite 3% 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 pm 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 (mass %) in the retained austenite divided by the average Mn content (mass %) in the ferrite and an average aspect ratio of the retained austenite is 3.0 or more,wherein the ratio of an average C content (7° by mass) in the retained austenite divided by an average C content (7° by mass) in the ferrite is equal to or greater than 3.0, wherein the ratio of the average C content (7° by mass) in the retained austenite divided by the average Mn content (7° by mass) in the retained austenite is less than 0.05, and wherein the diffusible hydrogen content in the steel is 0.3 ppm by mass or less, the method comprises hot rolling a steel slab having the chemical composition described in [1] or [2], then winding it at 300 °C or more and 750 °C or less, then cold rolling it, and then holding it at a temperature equal to or greater than a transformation temperature of 50 °C for 20 minutes or more. 1800 or less,then cooling to a cooling stop temperature equal to or lower than the martensitic transformation start temperature, then reheating to a reheat temperature equal to or higher than the transformation temperature Aci and equal to or lower than the transformation temperature Aci of +150 °C, subsequently holding the reheat temperature for 20 seconds more than and 1800 seconds less, then cooling to 100 °C or less, then holding a temperature above 100 °C and 400 °C or less for 10 seconds more, and subsequently cooling. [6] The method for manufacturing a high-strength steel sheet described in [5] further includes, after winding, holding 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 described in [5] or [6], the method further includes, after holding at a temperature equal to or higher than the transformation temperature Am and equal to or lower than the transformation temperature Aci of +150 °C for 20 seconds more and 1800 seconds less, performing a cooling, then performing a galvanizing treatment, and subsequently performing a cooling to 100 °C or less. [8] The method for manufacturing a high-strength steel sheet described in [7] further includes, after galvanizing treatment, performing an alloy treatment of the zinc coating 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 high yield strength (YR) of more than 0.70 and a tensile strength (TS) of 980 MPa or more, and that is excellent in terms of formability, specifically not only ductility but also the ability to be flanged and bent. 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 very 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 thus 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 obtain a sufficient area fraction of retained austenite, and adequate ductility may not be achieved. If excessive amounts of C are added to the steel, such that the C content exceeds 0.250%, the hardened martensite area fraction 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 quenched martensite may increase, and crack propagation may proceed unfavorably.Therefore, the ability to be drawn out by flanging 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. Based on the points 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 the Mn is concentrated, thus achieving adequate ductility. The actions described above are confirmed when the Mn content in the steel is 2.50% or more. However, if Mn is added to the steel in an excessive amount, such that the Mn content exceeds 8.00%, the area fraction of the quenched 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 quenched martensite may increase, and crack propagation may proceed unfavorably.Therefore, the ability to be drawn out by stretching may be degraded. Furthermore, the ease of conversion treatment and the quality of the coating may be degraded. From the perspectives described above, the Mn content is limited to 2.50% or more and 8.00% or less. The Mn content is preferably 3.10% or more and 6.00% or less, and more preferably 3.20% or more and 4.20% or less. 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 degraded. Furthermore, if the zinc coating alloying treatment is performed, the alloying rate may be reduced and the quality of the zinc coating may be compromised. 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%. The N content is preferably 0.0010% or more and 0.0070% or less. Aluminum (Al): 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 for 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 in steel. Ti increases ferrite strength, thus reducing the hardness difference between ferrite and the second hard phase (martensite or retained austenite), and consequently allows 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 excessively. In such a case, during a hole expansion test, the number of microvoids formed at the grain boundaries of the quenched 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 micro-voids formed at the grain boundaries of the quenched martensite may increase and crack propagation may proceed in a disadvantageous manner.Therefore, the flexural strength 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 (B) inhibits the formation and growth of ferrite at austenite grain boundaries. B increases the strength of ferrite, thus reducing the hardness difference between ferrite and the second hard phase (quenched martensite or retained austenite) and consequently enabling adequate drawability. If B is added to steel, the B content is preferably 0.0003% or more, and more preferably 0.0005% or more. However, if the B content exceeds 0.0050%, formability may be degraded. Therefore, if B is added to steel, the B content should be limited to 0.0050% or less. The B content is preferably 0.0030% or less. Ni: 1,000% or less Nickel (Ni) is an element that stabilizes retained austenite and is effective in achieving more suitable ductility. Ni also increases the strength of steel through solid solution strengthening. 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 quenched martensite increases excessively. In such a case, during a hole expansion test, the number of microvoids formed at the grain boundaries of the quenched martensite 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. If Cr and Mo are added to steel, the contents of each should preferably be 0.005% or more. 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 area fraction of hardened 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 hardened martensite may increase, and crack propagation may proceed unfavorably. Therefore, the drawability may be degraded. Consequently, if these elements are added to steel, the content of each should be limited to 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 hardened martensite increases excessively. In such a case, during a hole expansion test, the number of microvoids formed at the grain boundaries of the hardened martensite 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 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. This decarburization results from nitriding or oxidation of the steel sheet surface. Tin and antimony are effective in inhibiting this nitriding or oxidation, preventing a reduction in the area fraction of quenched martensite on the steel sheet surface and thus achieving a degree of strength and stability in mechanical properties. When tin and antimony are added to steel, the tin and antimony content is preferably 0.002% or higher. However, if either element is added to the steel in excessive amounts, such that the content exceeds 0.200%, toughness may be degraded. Therefore, if Sn and Sb are added to steel, the Sn and Sb content is limited to 0.200% or less. 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. Even if Ta is added to steel in excessive amounts, the precipitation stabilization effect can become saturated, and the cost of alloying elements also increases. 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 constituent other than those described above, i.e., the balance, 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—that is, ferrite materials that are relatively soft and ductile. The ferrite area fraction is preferably 40% or more and 75% or less. Area fraction of tempered martensite: 3.0% or more and 35% or less The quenched martensite area fraction must be 3.0% or more to achieve high local elongation, adequate drawability, adequate bendability, and a high yield ratio. It is necessary to limit the quenched martensite area fraction to 3.0% or more to achieve high local elongation, adequate drawability, adequate bendability, and a high creep coefficient. It is also necessary to limit the quenched martensite area fraction to 35% or less to achieve a tensile strength (TS) of 980 MPa or more. The quenched martensite area fraction is preferably 5.0% or more and 20% or less. The area fractions of the quenched 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 Cybernetics, 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 quenched martensite appears as a microstructure component that includes a white (as quenched) martensite portion and a gray internal structure formed within the martensite portion. 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 area fraction of retained austenite 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 pmo 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 to the total area fraction of retained austenite of 0.3 or more is an important condition of 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 prior to the flange forming step and thus improve 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 must be high, in addition to satisfying the total retained austenite area fraction, which is sufficiently high to achieve high ductility.The previous index is preferably 0.5 or higher. The upper limit for the previous aspect ratio is preferably 15.0 or lower. The lower limit for the previous minor axis length is preferably 0.05 pm or higher, 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 elements 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 iviA / a / zuzz / uu^o / 1 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 of the above ratio is not limited because the higher the average Mn content in the retained austenite, the greater the ductility, the above ratio is preferably 10.0 or less. 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 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. If the product 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 is less than 3.0, the probability of void formation during punching prior to flange forming increases significantly, and consequently, the drawability may be degraded. The above ratio is preferably 4.0 or higher.The upper limit for the above index is 20.0 or less. 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 greater is an extremely important condition that constitutes the present invention. 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 greater. The upper limit for the above ratio is 10.0 or less. The ratio between the average C content (% by mass) in the retained austenite and the average Mn content (% by mass) in the retained austenite is less than 0.05 Limiting the ratio of carbon content (mass %) in retained austenite to manganese content (mass %) in retained austenite to 0.05 or greater is an extremely important condition that constitutes the present invention. To achieve a high yield strength (YR), when manganese is concentrated in the retained austenite, including carbon and manganese in a proportion greater than carbon, the stability of the retained austenite is increased, and the yield strength is increased accordingly. The above ratio is preferably 0.02 or greater and 0.04 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 diffusible hydrogen content in steel is 0.3 ppm by mass or less Limiting the diffusible hydrogen content in steel to 0.3 ppm by mass or less is an important condition of the present invention. To achieve high local elongation and adequate flangability, it is necessary to limit the diffusible hydrogen content in the steel to 0.3 ppm by mass or less. The diffusible hydrogen content in the steel is preferably 0.2 ppm by mass or less. A test sample 30 mm long and 5 mm wide was taken from an annealed steel sheet. After removing the coating layer by grinding, the diffusible hydrogen content in the steel and the diffusible hydrogen release peak were measured. The release peak was measured by thermal desorption spectrometry (TDS) at a heating rate of 200 °C / hour. Hydrogen detected at 300 °C or less was considered diffusible hydrogen. 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 When heating the slab, the temperature to which it 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 its strength, the Ti and Nb-based precipitates formed during casting can be redissolved. It is also preferable to limit the heating temperature of the steel slab to 1100 °C or higher to reduce air bubbles, segregation, and similar issues 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 process to avoid problems occurring 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 reduction ratio of the roll is increased while the austenite is in a non-recrystallized state. In such a case, an abnormal texture develops. Consequently, the in-plane anisotropy of the final product increases significantly. This can degrade not only the uniformity of the material quality (stability of mechanical properties) but also the ductility. Furthermore, the aspect ratio of the retained austenite grains is reduced, and consequently, ductility and drawability can be degraded. Therefore, the delivery temperature for finishing in hot rolling is preferably between 750 °C or higher and 1000 °C or lower, and preferably between 800 °C or higher and 950 °C or lower. iviA / a / zuzz / uu^o / 1 Temperature at which winding is performed after hot rolling: 300 °C or more and 750 °C or less If the temperature at which rewinding 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 in the final annealed steel sheet decreases, and the flangability is consequently degraded. If the rewinding temperature after hot rolling is below 300 °C, the strength of the hot-rolled steel sheet increases. In such a case, the rolling load required for cold rolling may increase, and defects in the shape of the steel sheet may occur. This reduces productivity. Therefore, it is necessary to limit the rewinding temperature after hot rolling to 300 °C or higher and 750 °C or lower.The previous winding temperature is preferably 400 °C or more and 650 °C or less. 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 at a temperature equal to or lower than the Aci transformation temperature for more than 1800 s It is preferable to perform holding at a temperature equal to or lower than the transformation temperature Aci for more than 1800 s because it softens the steel sheet that will be cold-rolled in the subsequent step. If holding 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. If holding is performed at a temperature above the transformation temperature Aci, furthermore, massive cooled martensite and retained austenite form. The massive microstructure remains even after subsequent heat treatment. Consequently, the aspect ratio may be reduced and the drawability may be degraded. Even at a temperature equal to or lower than the transformation temperature Aci, if the holding time is iviA / a / zuzz / uu^o / 1 At 1800 or less, it is difficult to eliminate the stress left 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 the cold rolling with the aforementioned reduction ratio allows for the formation of a sufficiently recrystallized microstructure and improves the properties. 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.Furthermore, it becomes impossible to concentrate Mn on the surface to maintain the coating quality at a certain level. If the aforementioned maintenance is performed for more than 1800 s, the Mn concentration on the surface becomes saturated, the amounts of hardened martensite and retained austenite included in the surface of the steel sheet that has undergone the final annealing treatment increase, and the bendability may degrade accordingly. Cooling temperature at cooling stop equal to or lower than the starting temperature of the martensitic transformation If the cooling stop temperature is higher than the start temperature of the martensitic transformation, and the amount of cooled martensite to be transformed is small, retained austenite grain nuclei with a high aspect ratio do not form. In such a case, in the subsequent annealing step, retained austenite may form at the grain boundaries in a disadvantageous manner. IVIA / a / ZUZZ / UU^D / 1 This increases the amount of retained austenite grains with a low aspect ratio, making it impossible to form the intended microstructure. Furthermore, the amount of Mn concentrated in the retained austenite is reduced, and consequently, a high YR may not be achieved. The cooling stop temperature is preferably equal to or higher than the martensitic transformation start temperature (-250 °C) and equal to or lower than the martensitic transformation start temperature (-50 °C). Reheating to a reheating temperature equal to or higher than the transformation temperature Aci and equal to or lower than the transformation temperature Aci of +150 °C and subsequent holding at the reheating temperature for 20 seconds plus and 1800 seconds less. Performing the holding process 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 1800 seconds is an extremely important condition that constitutes the present invention. If the holding process 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 the strength is consequently reduced. Furthermore, carbon and manganese do not concentrate in the retained austenite to sufficient levels. Consequently, ductility and bendability are degraded. If the reheating temperature exceeds the transformation temperature Aci of +150 °C, the area fraction of martensite increases. In addition, the concentration of manganese in the austenite becomes saturated.Consequently, a sufficient area fraction of retained austenite does not form, and ductility degrades. Furthermore, massive retained austenite forms. This results in a reduction of the aspect ratio and degradation of drawability. The reheating temperature is preferably equal to or lower than the Aci transformation temperature of +100 °C. If holding is performed for more than 1800 s, the growth of austenite grains in the minor axis direction is accelerated. This results in a reduction of the aspect ratio and degradation of drawability. Additionally, carbon concentration in the retained austenite occurs. As a result, it becomes difficult to achieve the intended ratio of average carbon content (mass %) in retained austenite divided by average manganese content (mass %) in retained austenite and a high yield strength (YR).If galvanizing is to be carried out subsequently, cooling must be performed before the galvanizing process. The cooling stop temperature before galvanizing is preferably 350°C or higher and 550°C or lower. Galvanizing Treatment In the case of hot-dip galvanizing, after annealing (reheating to a reheat temperature equal to or higher than the transformation temperature Aci and equal to or lower than the transformation temperature Aci of +150 °C, and holding at the reheat temperature for 20 seconds above and 1800 seconds below), the steel sheet is cooled to a temperature equal to or higher than that of the galvanizing bath by gas blast cooling, furnace cooling, or similar means, and then immersed in a galvanizing bath at a temperature equal to or higher than 440 °C and equal to or lower than 500 °C for hot-dip galvanizing. The coating weight is then adjusted by gas blasting or similar means. For hot-dip galvanizing, a galvanizing bath with an aluminum content of 0.08% or more and 0.30% or less is preferably used.Instead of hot-dip galvanizing treatment, electrogalvanizing or similar treatment can be used. 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. Cooling to 100 °C less Cooling to 100°C or below allows a sufficient amount of tempered martensite to form, enabling the achievement of a specific tensile strength after the final annealing treatment. Furthermore, it is preferable to perform cooling to approximately 20°C or higher and then to approximately 50°C or lower before the subsequent holding step of holding at 100°C or higher and then at 400°C or lower for 10 minutes or more. If galvanizing or zinc plating is to be performed, it is carried out before this holding step, i.e., before the cooling step to 100°C or lower. Maintain at over 100°C and 400°C or less for 10 minutes or more Performing the holding process at temperatures above 100°C and below 400°C for 10 seconds or less as a final heat treatment (final annealing) is important in the present invention. If the holding process is performed at 100°C or less, or for less than 10 seconds, a sufficient area fraction of quenched martensite does not form. Furthermore, the diffusible hydrogen contained in the steel does not detach from the steel sheet. This results in a degradation of the steel's drawability and bendability. If the holding process is performed at temperatures above 400°C, a sufficient area fraction of retained austenite does not form due to the decomposition of the retained austenite. Consequently, the ductility of the steel is degraded. 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 aforementioned high-strength steel sheet and high-strength galvanized steel sheet, which includes the high-strength steel sheet and a galvanized coating on the surface, can undergo 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 rolling reduction ratio is less than 0.1%, the benefits are minimal, and the reduction ratio is difficult to control. 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 reduced. 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 roll reduction ratio can be carried out in one or more passes. Furthermore, 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 mm to 1.8 mm under the conditions described in Table 2. A galvanizing treatment was then carried out to produce hot-dip galvanized (Gl) steel sheets. The hot-dip galvanized steel sheets were then 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 3, 4, and 5. Table 1 Steel Type Chemical Composition (Mass) Ms Temperature (=C) Ac Transformation Temperature (=Ci) Ac Transformation Temperature: (=C Observation is C Si Mn PSN Al Ti Nb VWB Ni Cr M0 Cu Snsb Ta Ca Mg Zr RE MA 0 16 6 0 5 6 3 5 3 0. 02 2 00 02 1 00 03 5 0. 03 1 0 04 5 - 352 656 771 Invention Steel B 0 18 6 0 7 6 3 2 3 0. 02 7 0 0 02 5 0 0 04 2 0. 04 6 0 03 9 - 357 666 785 Invention Steel C 0 17 8 1 8 1 3 6 0 0. 01 L O 0 0 01 9 0 0 02 2 0. 03 4 0 03 2 - 345 667 817 Invention steel D 0 24 3 0 9 8 3 3 0 0. 02 L O 0 0 03 0 00 02 6 0. 05 9 - 335 665 767 Invention steel E 0 05 0 1 0 1 4 0 9 0. 03 1 0 0 02 4 0 0 02 7 0. 03 2 - 370 647 794 Invention steel F 0 18 2 2 8 9 3 9 8 0. 02 6 0 0 02 0 0 0 02 5 0. 03 2 0 04 3 - 328 668 858 Invention steel G 0 19 9 0 6 2 3 5 3 0. 03 0 0 0 01 9 00 03 4 0. 04 2 0 05 0 - 340 656 770 Invention steel H 0 08 3 1 0 0 5 1 3 0. 02 7 0 0 02 7 0 0 03 2 0. 04 6 - 317 617 752 Invention steel 1 0 17 9 1 5 0 3 7 8 0.02 0 0 0 02 1 0 0 02 6 0. 03 6 - 337 659 785 Invention steel J 0 15 8 0 2 1 3 5 0 0. 03 0 0 0 02 4 0 0 03 8 0. 03 1 0 04 6 - 356 653 758 Invention steel K 0 12 6 0 3 7 5 9 4 0. 02 6 0 0 02 5 00 03 2 0. 03 1 0 05 0 - 269 587 703 Invention steel. L 0 19 1 0 4 6 3 1 4 0. 02 3 0 0 02 6 0 0 02 7 0. 03 3 - 359 665 754 Invention steel M 0 15 4 0 6 2 4 2 0 0. 02 3 0 0 02 4 0 0 03 1 0 03 6 - 329 638 739 Invention steel N 0 17 2 0 4 9 3 2 1 0. 03 0 0 0 02 1 0 0 04 3 0 03 7 0. 04 0 - 363 664 775 Invention steel 0 0 16 1 0 8 2 3 4 3 0. 01 4 0 0 01 9 0 0 03 5 0. 71 1 0. 03 9 - 378 661 920 Invention steel P 0 15 7 0 5 9 3 5 4 0. 01 9 0 0 02 7 0 0 02 9 1. 17 9 0. 04 3 - 389 656 100 3 Invention steel Q 0 20 1 0 3 5 3 5 0 0. 02 6 0 0 02 7 0 0 04 3 0 22 2 - 346 654 774 Invention steel R 0. 02 2 0 5 1 3 5 2 0. 02 0 0 0 02 8 0 0 03 9 0.03 7 0.04 2 - 403 658 821 Comparative steel S 0 20 6 l'j lo lt- 3 5 0 0.02 9 0 0 02 5 0 0 03 5 0.03 1 - 339 694 902 Comparative steel T 0 18 5 0 3 4 8 4 4 0.02 5 0 0 02 4 0 0 02 6 0 03 4 - 149 515 592 Comparative steel u 0 15 7 1 0 9 HO 0.02 1 0 0 02 2 0 0 02 8 0 03 0 0.04 2 - 408 699 836 Comparative steel V 0 16 5 0 5 9 2 5 7 0. 01 9 0 0 01 8 0 0 04 1 0.04 2 L 24 9 - 391 683 885 Comparative steel w 0 15 5 0 7 5 3 2 3 0. 01 9 0 0 02 4 0 0 02 9 0. 05 2 - 0. 06 3 - 368 666 777 Invention steel X 0 16 1 0 6 9 4 5 1 0. 03 0 0 0 02 4 0 0 04 1 0 04 6 0. 01 1 0. 03 9 - 315 630 738 Invention steel Y 0 12 0 1 1 4 3 5 9 0. 03 2 0 0 02 3 0 0 02 6 0. 04 5 0.09 0 - 0.06 1 - - - - - - - - - - - - - 364 661 835 Invention steel z 0 10 2 1 2 0 4 0 7 0.02 ü 0 0 02 7 0 0 03 1 0.04 4 0.9 - 353 649 786 Invention steel AA 0 14 9 0 3 5 3 4 6 0 03 0 0 0 02 3 0 0 04 4 0 04 0 0 01 9 - - - 00 02 0 - - - - - - - - - - - 361 656 759 Invention steel. 0 0 5 0. 0 0 0 0 0. 0. 0. Invention steel AB 18 8 6 7 9 6 02 5 02 2 04 0 01 5 01 1 - - - - 04 2 - - - - - - - - - - 246 588 680 0 0 6 0. 0 0 0 0 0. 0. 0. Invention steel AC 09 2 5 2 3 9 01 9 02 7 03 8 05 ü 07 2 - - - - - 09 O - - - - - - - - 262 578 722 0 0 3 0. 0 0 0 0 0. 0. 0. Invention steel AD 14 9 8 2 6 4 02 1 01 9 03 7 03 2 04 9 - - - - - 52 7 - - - - - - - - - 343 663 791 0 1 3 0. 0 0 0 0 0. 0. 0. Invention Steel AE 09 9 4 4 0 8 03 0 02 7 03 1 03 3 02 5 - - - - - - 06 1 - - - - - - - - 393 679 837 AF 0 10 8 0 5 0 3 5 9 0. 02 4 0 0 02 4 0 0 02 6 0. 04 1 0. 30 1 368 653 762 Invention Steel 0 0 3 0. 0 0 0 0 0. 0. 0. Invention steel AG 11 8 5 5 1 7 02 5 02 1 03 2 03 4 02 5 - 10 5 383 666 787 0 0 3 0. 0 0 0 0 0. 0. 0. Invention steel AH 16 1 4 2 2 5 02 1 02 2 02 7 03 2 03 5 - - - - - - - - - 05 0 - - - - 365 662 770 0 0 3 0. 0 0 0 0 0. 0. Invention steel Al 13 5 7 0 5 8 01 9 02 0 02 9 03 0 - 00 7 360 656 766 0 0 3 0. 0 0 0 0 0. 0. 0.Invention steel AJ 20 0 4 0 0 1 03 0 02 9 02 7 02 Λ Ü - 04 0 - - - - - - - 00 7 - - - - - - 360 668 753 AK 0 21 1 0 2 3 3 6 9 0. 02 3 0 0 02 7 0 0 03 9 0. 03 3 - 0. 03 2 - - - - - - - - - 0. 00 8 - - - - 330 647 723 Invention steel 0 0 3 0. 0 0 0 0 0. 00 Invention steel AL 21 0 9 7 9 6 02 5 02 5 04 1 04 0 - 03 4 319 647 750 AM 0 19 5 1 2 5 3 7 9 0. 02 1 0 0 02 3 0 0 03 5 0. 03 5 - - - - - - - - - - - - - 00 02 5 - - 331 656 770 Invention Steel 0 0 3 0. 0 0 0 0 0. 0. 00 Invention Steel AN 24 1 0 4 0 1 02 6 02 2 02 8 03 9 00 7 - 03 0 - 346 663 732 0 0 6 0. 0 0 0 0 0. 0 0 Invention Steel AO 07 8 0 6 1 1 02 2 02 8 03 7 04 2 - 03 1 280 579 681. Underlined portion: Outside the scope of the present invention, -: Content at the level of incidental impurities The starting temperature of the martensitic transformation, i.e., the Ms temperature, and the transformation temperatures Aci and Acsse were calculated using the following formulas. Onset temperature of martensitic transformation: Temperature Ms (°C) = 550 - 350 χ (% of C) - 40 χ (% of Μη) - 10 χ (% of Cu) -17 x (% of Ni) - 20 x (% of Cr) - 10 x (% of Mo) - 5% of V - 55 χ 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 Acs (°C) = 910 - 203^ (% 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 Mo) + 10 χ (% of Mo) 400 χ (% Ti) + 200 χ (% Al) where: (% C), (% Si), (% Mn), (% Ni), (% Cu), (% Cr), (% Mo), (% V), (% Ti), (% V), (% W) and (% Al) represent each of the (%) contents by mass of the element. Table 2 N Tip 0 of steel Final rolling delivery temperature (=Ci Coil temperature rC) Heat treatment of hot rolled steel sheet Cold rolling reduction coefficient (%) Annealing treatment of cold rolled steel sheet Alloy temperature (=C) Final annealing treatment Type Heat treatment temperature (=Ci Heat treatment time 0 (if Heat treatment temperature l=C) Heat treatment time 0 (if Cooling stop temperature (=0) Reheating temperature CCi Reheating temperature holding time (if Heat treatment temperature (=C) Heat treatment time 0 (if 1 A 900 500 560 32400 61 1 850 150 130 675 180 110 72000 CR 2 A 890 440 510 10800 52 9 800 120 160 720 270 170 1200 Gl 3 A 860 490 590 10800 56 3 780 130 160 670 150 150 60000 Gl 4 A 910 460 550 14400 64 7 815 120 130 700 170 525 200 54000 GA 5 At 930 500 60010800 62 5 610 90 210 690 180 120 4200 Gl 6 A 905 530 590 9000 61 1 820 15 160 685 250 540 300 12000 G A 7 A 920 480 590 13000 588 870 120 410 680 130 530 310 14400 G A 8 A 790 540 570 32400 56 3 845 100 170 860 370 210 60000 C R 9 A 810 530 64 7 890 120 125 610 250 140 24400 Gl 1 0 A 860 440 580 18000 61 1 910 240 80 690 3600 505 320 7200 G A 1 1 A 875 480 510 9000 61 1 810 130 130 705 5 200 9000 C R 1 2 A 870 530 46 2 830 250 190 690 490 530 130 21600 GA 1 3 B 900 480 560 21600 548 825 360 170 705 550 170 25200 Gl C 900 510 570 7200 52 9 860 140 315 690 160 490 200 28300 GA 1 5 A 700 560 550 36000 47 1 860 60 250 690 265 500 230 600 GA 1 6 A 890 850 580 10800 56 5 860 210 140 730 240 490 220 14400 GA 1 7 A 890 610 730 18000 588 900 180 110 685 260 520 160 54400 GA 1 δ A 870 555 540 18000 50 0 895 90 130 680 370 35 13000 CR 1 9 A 900 540 560 21600 55 6 890 120 160 680 130 530 560 360 GA 2 0 A 895 570 565 24000 53 3 925 140 110 700 130 160 5 Gl 2 1 D 910 560 580 10800 58 8 890 210 130 715 290 515 380 60 GA fxj hJ | E 890 530 570 21600 588 920 420 170 690 540 300 23400 Gl 2 3 F 905 590 600 18000 57 1 840 75 190 710 235 170 36000 CR 4 G 870 600 580 9000 57 1 800 120 190 680 140 510 300 13000 GA 2 5 H 875 510 505 21600 53 3 840 120 120 715 130 560 400 14400 GA 2 6 I 900 560 530 10800 50 0 770 150 160 690 170 530 230 57600 GA 2 7 J 890 520 520 18000 529 790 180 155 770590 210 18000 C R 2 δ K 860 560 530 32400 48 6 790 90 90 690 110 250 7200 Gl 2 9 L 910 590 560 23400 462 770 240 150 700 290 380 2400 Gl 3 0 M 870 620 580 36000 62 5 820 150 140 710 180 550 345 10300 G A 3 1 N 860 590 520 7200 62 5 800 140 130 670 170 265 36000 C R 3 2 0 880 500 520 10800 58 8 910 75 65 680 120 500 260 12000 G A 3 3 P 870 480 52 0 980 120 200 730 140 530 330 12000 G A 3 4 Q 880 530 610 36000 56 3 810 300 130 685 400 360 72000 C R 3 5 R 910 540 588 850 90 160 690 150 535 210 10800 G A 3 6 S 900 610 540 21600 62 5 900 150 110 695 190 480 310 4300 G A 3 7 S 875 550 580 10800 61 1 920 95 150 705 140 110 96000 C R 3 8 T 890 600 500 13000 62 5 830 160 30 640 200 160 18800 0 C R 3 9 u 915 520 520 10800 57 1 875 150 170 715 130 320 5400 Gl 4 0 y 850 480 510 18000 64 7 900 60 170 740 170 390 240 Gl 4 1 w 880 580 56 3 880 130 200 700 210 540 130 96000 G A 4 2 X 870 600 560 21600 50 0 800 360 200 715 520 300 24000 C R 4 3 Y 900 560 500 36000 46 2 850 150 140 710 200 540 240 9000 G A 4 4 z 870 540 560 14400 52 9 840 140 240 700 160 250 9600 C R 4 5 AA 905 580 47 1 850 300 90 720 420 510 260 14400 G A 4 6 AB 910 590 530 14400 55 6 840 180 200 640 250 525 280 72000 G A 4 7 AC 900 540 530 13000 56 3 820 130 180 620 1200 330 1200 Gl 4 8 AD 895 530 540 21600 53 3 830 120 190 710 220 525 125 60000 G A 4 9 AE 880 520 530 23400 58 8 790 120 65 720 170 110 60000 Gl 5 0 AF 860 510 510 23400 64 7 850 120 170 630 120 150 54000 Gl 5 1 AG 910 500 590 9000 62 5 850 150 140 680 200 520 2604200 GA 5 2 AH 920 570 520 28800 56 3 870 160 140 700 210 540 150 12000 GA 5 3 Al 845 590 53 8 870 330 130 760 190 510 280 72000 GA 5 4 AJ 920 560 520 24400 56 3 900 90 140 690 130 200 4300 Gl 5 5 AK 910 540 540 13600 56 3 860 30 220 710 1350 500 160 96000 GA 5 6 AL 840 540 530 14400 56 3 810 120 180 730 140 120 18300 0 CR 5 7 AM 880 530 500 10800 64 7 840 150 170 750 195 540 190 9600 GA 5 8 AN 840 520 46 7 850 95 80 715 190 220 24000 CR 5 9 AO 860 500 560 31200 50 0 840 300 90 635 300 530 150 36000 GA 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 iviA / a / zuzz / uu^o / 1 No. Steel Type Thickness (mm) Area Fraction of F (%) Area Fraction of TM (%) Area Fraction of RA (%) Area Fraction of AR with aspect ratio equal to or greater than 2.0 and minor shaft length equal to or less than 1 pm / Total Area Fraction of RA Average Mn Content in RA (% by mass) Average Mn Content in F (% by mass) Average Mn Content in RA / Average Mn Content in F Average Aspect Ratio of RA 1 A 1.4 55.4 18 7 20.5 0 342 7 46 2.11 3 54 5.40 2 A 1.6 53.7 20 0 19 0 0 388 7 20 2.85 2 53 4.34 3 A 1.4 54.2 20.3 22.2 0 492 788 2.78 2 83 4.88 4 A 1.2 49 6 23.4 18 0 0 544 7 43 2.72 2 73 5 13 5 A 1.2 55.2 25.4 11.6 0 197 688 3.02 228 1.04 6 A 1.4 29.4 29.0 10.9 0 413 763 2.68 285 3.99 7 A 1.4 54.2 22 8 13 4 0 120 6 32 3.41 1 85 1.00 3 A 1.4 51 3 38 3 74 0 211 7 45 2.81 2 65 4.54 9 A 1.2 75.2 2 4 2.5 0 400 4 13 2.98 1 39 3.99 10 A 1.4 51 4 23 3 21 4 0 144 6 11 2 91 2 10 3.46 11 A 1.4 74.1 2.3 58 0 325 4 10 3.22 1 27 3.70 12 A 1.4 52.3 18.4 19.7 0 553 9 12 2.87 3 17 5.64 13 B 1.4 54.0 17.2 19 8 0 301 8 11 2.83 2 87 7.11 14 C 1.6 54.8 18.9 20.5 0 591 8 13 1.73 4 70 5 15 15 A 1.8 50.7 20.4 22.2 0 611 7 16 1.23 5 82 4.22 16 A 1.0 51 8 18.7 24 7 0 156 7 76 1 81 429 2.22 17 A 1.4 55 8 23.5 18.0 0 234 7 42 3.21 2 31 3.13 18 A 1.4 55 8 14 191 0 344 7 60 2.84 2 68 3.44 19 A 1.6 58 3 30 1 4.8 0 399 8 46 3.10 2 73 4.55 20 A 1.4 57.9 1.9 21.3 0 513 8 01 2.55 3 14 6.14 21 D 1.4 60.6 18.9 18.6 0 630 8 07 2.71 2 98 5.37 22 E 1.4 53.0 14.6 23.5 0 484 7 66 2.78 2 76 6.21 23 F 1.2 54.7 22.4 19.4 0 418 8 17 2.10 3 89 5.14 24 G 1.2 46.4 28.9 15 3 0 359 8 78 2.45 3 58 3.89 25 H 1.4 55.4 20 5 14.6 0 498 10 09 3.11 3 24 3.99 26 I 1.4 53.3 24.5 15.2 0 701 7 89 2.88 2 74 4.58. 27 J 1.6 53 1 18.6 25.9 0 362 5 67 2.24 2 53 5.39 28 K 1.8 49.6 21.9 24.5 0 743 10 94 4.44 2 46 4.67 29 L 1.4 58.5 20.3 15.2 0 303 628 2.87 2 19 6.34 30 M 1.2 51.3 21.3 23.1 0 575 8 25 2.47 3 34 5.47 31 N 1.2 51.7 23.3 16.9 0 400 8 99 2.91 3 09 8.30 32 0 1.4 52 7 19 9 199 0 330 7 75 3.22 2 41 6.45 33 P 1.2 50.2 19.0 25.8 0 410 7 91 1.92 4 12 3.30 34 Q 1.4 49.4 28.0 14.1 0 350 5 95 1.07 5 56 4.98 35 R 1.4 79.5 2.2 7.2 0 322 6 04 2.45 2 47 2.99 36 S 1.2 51.6 14.6 25.2 0 666 7 17 2.66 2 70 5.11 37 S 1.4 53 2 14.6 26.1 0 645 6 86 2.96 2 32 4.47 38 T 1.2 40.2 39.3 20.3 0 306 9 96 6.12 1 63 6.49 39 u 1.2 58.8 28.4 6 3 0 401 3 62 2.61 1 39 3.98 40 y 1.2 45.5 36.4 16.4 0 470 4 94 2.12 2 33 5.44 41 w 1.4 55.8 17.2 25.4 0 444 6 33 3.10 2 04 4.15 42 X 1.4 56.7 19.0 19.4 0 312 7 13 2.99 2 38 5.32 43 Y 1.4 52.6 19.6 24.6 0 304 6 34 2.40 2 64 4.28 44 z 1.6 57.9 15.5 20.3 0 396 6 58 3.02 2 18 5.13 45 AA 1.8 51.2 22.2 23.4 0 409 5 27 2.55 2 07 6.20 46 AB 1.6 49.8 13.3 29.5 0 459 10 96 2.92 3 75 6.09 47 AC 1.4 50.1 21.4 22.0 0 403 10 44 4.22 2 47 4.53 48 AD 1.4 50.2 22.3 22.2 0 472 7 23 3.11 2 32 4.91 49 AE 1.4 54.0 17.8 23.0 0 378 8 34 2.09 3 99 5.33 50 AF 1.2 53 9 17 5 22 6 0 583 7 75 1 92 4 04 6.04 51 AG 1.2 52.0 21 3 21.2 0 421 9 00 3.05 2 95 5.89 52 AH 1.4 49.0 32.9 13 6 0 581 5 99 2.95 2 03 4.07 53 Al 1.2 48.6 29.9 16 0 0 354 7 12 2.75 2 59 5.11 54 AJ 1.4 51.2 20.3 20.4 0 449 6 60 3.31 1 99 4.81 55 AK 1.4 49.6 20.4 20.2 0 328 7 63 3.27 2 33 5.41 56 AL 1.4 50.2 20.0 20.4 0 309 7 14 2.59 2 76 4.14 57 AM 1.2 53.4 20 1 24.3 0 383 6 60 3.05 2 16 5.54 58 AN 1.6 52.7 20.3 20.7 0 335 8 77 1.70 5 16 4.52 59 AO 1.4 52 1 20.2 20.8 0 564 12 04 4.50 2 68 5.01. Underlined portion: Outside the scope of the present invention F: Ferrite, TM: Tempered martensite, RA: Retained austenite The tensile test was performed in accordance with JIS Z 2241 (2011) using a JIS test specimen no. 5 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, the tensile strength (TS), elongation (EL), yield strength (YS), and yield ratio (YR) were measured, where YR is the ratio of YS to TS. In the present invention, a "Good" rating was given for the mechanical properties when the YR was greater than 0.70. A "Good" rating was also given for the mechanical properties 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 Df 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 Z 2248 (1996), using a 30 mm wide and 100 mm long bendability test specimen 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. Table 4 No. Average Mn content in AR / Average Mn content in F Average aspect ratio of RA Average C content in RA (i% by mass) Average C content in F (% by mass) Average C content in RA' Average C content in F Average C content in ARj' Average Mn content in RA Equilibrium microstructure YS (MPai) YR TS (MPai) EL (%) 1 1909 0.32 0.05 6 40 0.04 MBFP6 722 0.72 996 22 5 2 10 96 0.31 0.04 7 75 0.04 M BF.P5 743 0.73 1022 25 1 3 13 83 0.35 0.06 5 83 0.04 M BF.P6 825 0.82 1006 25 6 4 1401 0.31 0.08 3 88 0.03 MBF.P6 829 0.75 1099 20 9 5 2.37 0.36 0.06 6 00 0.04 MBF.P9 789 0.77 1022 17 3 6 11 36 0.19 0.04 4 75 0.02 M BF P 5 880 0.83 1055 14 2 7 1.85 0.43 0.05 8 60 0.07 M BF P 5 646 0.63 1032 16 8 Λ 0 12 04 0.33 0.09 3 67 0.04 M BFP6 822 0.78 1052 18 3 9 5 53 0.16 0.06 2 67 0.04 M BF P6 702 0.72 975 14 1 10 7.26 0.35 0.06 6 11 0.06 MBF.P6 665 0.62 1079 22 4 11 4.71 0.17 0.09 1 89 0.04 M BF P5 654 0.71 919 14 5 12 17 90 0.29 0.03 8 83 0.03 M BF P 5 784 0.75 1052 26 5 13 20 38 0.36 0.08 4 68 0.04 M BF P 5 696 0.71 982 22 9 14 2421 0.24 0.07 3 43 0.03 M BF P5 811 0.79 1031 23 5 15 24 57 0.30 0.10 3 00 0.04 M BF P6 820 0.78 1050 20 1 16 9.52 0.25 0.05 5 00 0.03 M BF P5 713 0.72 996 27 4 17 7.24 0.33 0.06 5 39 0.04 M BF.P6 913 0.81 1130 25 4 18 9.21 0.30 0.05 6 00 0.04 M BF P 9 811 0.77 1058 26 3 19 12 42 0.28 0.04 7 00 0.03 MBF.P9 723 0.72 999 12 4 20 1929 0.35 0.05 7 00 0.04 MBFP9 741 0.72 1025 25 0 21 15 99 0.34 0.07 4 55 0.04 MBFP9 944 0.90 1046 20 6 22 17 14 0.30 0.04 7 50 0.04 M BF.P9 710 0.72 983 22 3 23 20 00 0.30 0.09 3 33 0.04 M BF.P9 883 0.76 1161 22 1 24 13 94 0.22 0.05 4 40 0.03 MBF.P9 832 0.77 1078 23 7 25 12 95 0.24 0.04 6 00 0.02 MBF.P9 751 0.75 1001 20 9 26 12 55 0.30 0.07 4 29 0.04 MBFP9 763 0.74 1035 24 1 27 13 65 0.23 0.06 3 83 0.04 MBFP9 712 0.72 988 26 2 28 11 51 0.27 0.08 3 38 0.02 M BF.P9 742 0.74 1004 28 1 29 13 87 0.28 0.05 5 60 0.04 MBF.P9 881 0.81 1084 20 9 30 1824 0.35 0.11 3 13 0.04 MBF.P9 852 0.81 1054 23 9 31 25 64 0.20 0.04 5 00 0.02 MBF.P9 821 0.80 1022 22 1 32 15 52 0.32 0.07 4 57 0.04 MBFP9 819 0.82 993 21 5. 33 1361 0.35 0.05 7 00 0.04 MBF.P5 760 0.74 1024 23 0 34 27 68 0.25 0.06 4 17 0.04 M BF.P6 858 0.81 1064 21 3 35 7.37 0.18 0.02 9 00 0.03 MBF.P9 795 0.94 844 15 5 36 13 77 0.32 0.04 8 00 0.04 MBF.P9 738 0.75 989 13 8 37 10 36 0.25 0.06 4 17 0.04 M BF.P9 776 0.79 985 14 2 38 10 56 0.42 0.07 6 00 0.04 M BFP5 800 0.73 1102 30 3 39 5.52 0.41 0.05 8 20 0.11 M BFP6 647 0.63 1028 13 9 40 1266 0.22 0.05 4 40 0.04 M BF.P6 864 0.84 1023 24 7 41 8.47 0.25 0.06 4 17 0.04 M BF P6 721 0.71 1016 26 3 42 1269 0.32 0.05 6 40 0 04 M BF.P6 925 0.82 1127 23 2 43 11 31 0.28 0.06 4 67 0.04 MBF.P5 736 0.73 1002 22 9 44 11 18 0.24 0.07 3 43 0.04 M BF.P9 755 0.76 988 21 2 45 1282 0.20 0.05 4 00 0.04 M BF.P6 832 0.78 1073 23 2 46 2286 0.15 0.04 3 75 0 01 M BF.P3 779 0.79 983 30 4 47 11 21 0.21 0.04 5 25 0.02 MBF.P9 770 0.76 1015 20 7 48 11 41 0 31 0.06 5 17 0 04 M BF.P6 779 0.72 1085 21 2 49 21 27 0.29 0.04 7 25 0.03 M BF P5 724 0.72 1002 21 5 50 24 38 0.33 0.06 5 32 0.04 M BFP5 889 0.89 997 24 1 51 17 38 0.26 0.08 3 25 0.03 M BF.P9 865 0.78 1105 24 3 52 8.26 0.24 0.04 6 00 0.04 MBF.P6 900 0.83 1090 21 1 53 13 23 0.31 0.06 5 17 0.04 M BF.P3 744 0.72 1033 22 4 54 9 59 0 29 0.09 3 22 0 04 M BF P 3 793 0.79 998 26 8 55 12 62 0.34 0.10 3 27 0.04 M BF.P3 817 0.83 981 23 2 56 11 41 0.31 0.07 4 43 0.04 MBF.P5 754 0.73 1027 22 4 57 11 99 0.28 0.09 3 11 0.04 MBF.P3 769 0.77 995 26 6 58 23 32 0.36 0.07 5 14 0.04 M BF.P3 749 0.76 992 20 9 59 13 40 0.28 0.03 9 33 0.02 MBF.P5 821 0.83 986 22 1. Underlined portion: Outside the scope of the present invention F: Ferrite, M: Tempered martensite, RA: Retained austenite, BF: Bainitic ferrite P: Pearlite, Θ: Carbides (e.g., cementite) Table 5 No. Λ R (mm) RA Diffusive hydrogen content in steel (mass ppm) Remarks 1 27 2.6 1 9 0.03 Example of invention 2 25 34 2 1 0.06 Example of invention 3 27 30 2 1 0.16 Example of invention 4 25 2 2 1 8 0.13 Example of invention 5 14 2.4 20 0.05 Comparative Example 6 32 30 2 1 0.04 Comparative Example 7 14 26 1 9 0.04 Comparative Example 3 11 2.4 1 7 0.07 Comparative Example 9 30 32 2 7 0.02 Comparative Example 10 13 32 2 3 0.05 Comparative Example 11 27 38 2 7 0.09 Comparative Example 12 23 32 23 0.16 Invention Example 13 27 16 1 1 0.13 Invention Example 14 29 16 1 0 0.23 Invention Example 15 28 2 8 1 6 0.11 Invention Example 16 11 20 20 0.20 Example Comparative Example 17 10 26 1 9 0.17 Comparative Example 18 13 4 0 29 0.42 Comparative Example 19 24 2 0 1 3 0.09 Comparative Example 20 13 3 8 2 7 0.39 Comparative Example 21 21 2.4 1 7 0.22 Invention Example 22 30 32 23 0.05 Invention Example 23 25 12 1 0 0.17 Invention Example 24 25 2 6 2 1 0.03 Invention Example 25 30 0 6 0 5 0.11 Invention Example 26 24 2 8 20 0.24 Invention Example 27 22 26 1 6 0.23 Invention Example 28 28 32 1 3 0.21 Example of invention 29 26 23 1 7 0.24 Example of invention 30 27 2.2 1 8 0.24 Example of invention 31 26 28 23 0.08 Example of invention 32 26 30 2 1 0.17 Example of invention 33 26 2.2 1 8 0.02 Example of invention 34 22 2 6 1 9 0.06 Example of invention 35 29 2.2 1 6 0.01 Comparative example 36 31 48 40 0.15 Comparative example 37 25 49 35 0.11 Comparative example 38 9 2 6 2 1 0.30 Comparative example 39 19 2 2 1 8 0.06 Comparative example 40 11 24 20 0.12 Comparative example 41 24 20 1 4 0.05 Example of invention 42 27 2 4 1 7 0.17 Example of invention 43 25 2 6 1 9 0.21 Example of invention 44 25 26 1 6 0.23 Example of invention 45 32 2.2 1 2 0.11 Example of invention 46 24 26 1 6 0.25 Example of invention 47 24 3 0 2 2 0.19 Example of invention 48 23 2.2 1 6 0.04 Example of invention. 49 26 26 1 9 0.22 Example of invention 50 27 16 1 3 0.29 Example of invention 51 26 12 1 0 0.18 Example of invention 52 25 2 8 2 0 0.09 Example of invention 53 24 24 20 0.22 Example of invention 54 28 2 8 2 0 0.16 Example of invention 55 27 2 4 1 7 0.21 Example of invention 56 28 22 1 6 0.17 Example of invention 57 28 2 8 2 3 0.08 Example of invention 58 30 22 1 4 0.29 Example of invention 59 26 2 8 2 0 0.10 Example of invention IVIA / a / ¿U¿¿ / UU40 / 1 Underlined portion: Outside the scope of the present invention Industrial Applicability According to the present invention, a high-strength steel sheet can be produced with a yield strength (YR) greater than 0.70, 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, quenched martensite: 3.0% 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 pm 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 more, 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, 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 less than 0.05, and wherein a diffusible hydrogen content in the steel is 0.3 ppm by mass or less.

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 comprising a steel microstructure that includes, by area, ferrite 30% or more and 80% or less, quenched martensite 3% or more and 35% or less, and retained austenite 8% or more, wherein the 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 pm or less, divided by a total area fraction of the retained austenite is 0.3 or more, wherein the 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 the product of the ratio of the average Mn content (mass %) in the retained austenite divided by the average Mn content (mass %) in the ferrite and an average aspect ratio of the retained austenite is 3.0 or more,wherein the ratio of an average C content (% by mass) in the retained austenite divided by an average C content (% by mass) in the ferrite is equal to or greater than 3.0, 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 less than 0.05, and wherein the diffusible hydrogen content in the steel is 0.3 ppm by mass or less, the method comprises hot rolling a steel slab having the chemical composition according to claim 1 or 2, then winding it at 300 °C or more and 750 °C or less, then cold rolling it, and then holding it at a temperature equal to or greater than a transformation temperature of 50 °C for 20 seconds or more and 1800 seconds or less.then cooling to a cooling stop temperature equal to or lower than the martensitic transformation start temperature, then reheating to a reheating temperature equal to or higher than the transformation temperature Aci and equal to or lower than the transformation temperature Aci of +150 °C, subsequently holding the reheating temperature for 20 °C or more and 1800 °C or less, then cooling to 100 °C or less, then holding a temperature above 100 °C and 400 °C or less for 10 °C or more, and subsequently 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 according to claim 5 or 6, the method further comprising, after holding 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 more and 1800 seconds less, performing a cooling, then performing a galvanizing treatment, and subsequently performing a cooling to 100 °C or less.

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 of the zinc coating at 450 °C or more and 600 °C or less.