Source: http://patents.com/us-20170096723.html
Timestamp: 2019-04-22 13:04:23+00:00

Document:
Provided is a high strength cold rolled steel sheet having a tensile strength of 980 MPa or higher and excellent ductility and bendability. The high strength cold rolled steel sheet has a specified component composition. The structures at a position of 1/4 sheet thickness in the steel sheet has certain area ratio of ferrite relative to the entire structure, certain area ratio of a mixed structure of fresh martensite and retained austenite relative to the entire structure, certain region in which the concentration of Mn is enriched, certain standard deviation of the fraction of a region in which the concentration of Mn is enriched, and certain concentration of Mn in a ferrite phase.
1. A high strength cold rolled steel sheet having a tensile strength of 980 MPa or higher and excellent ductility and bendability, the steel sheet satisfying a component composition comprising in percent by mass: C: 0.10% to 0,30%, Si: 1.2% to 3%; Mn: 0.5% to 3.0%; P: more than 0% to 0.1% or less; S: more than 0% to 0.05% or less; Al: 0.005% to 0.2%; N: more than 0% to 0.01% or less; and O: more than 0% to 0.01% or less, the balance being iron and inevitable impurities, and structures at a position of 1/4 sheet thickness in the steel sheet satisfying all of (1) to (5) below: (1) the area ratio of ferrite relative to the entire structure is from 5% or more to less than 50%, and the balance is a hard phase, when observed under a scanning electron microscope; (2) the area ratio of a mixed structure of fresh martensite and retained austenite relative to the entire structure is more than 0% to 30% or less, when observed under an optical microscope after LePera etching; (3) a region in which the concentration of Mn is enriched to 1.2 times or more the concentration of Mn in the steel sheet is present at 5 area% or more, when analyzed with an electron probe microanalyzer; (4) when the fraction of a region in which the concentration of Mn is enriched to 1.2 times or more the concentration of Mn in the steel sheet is measured in 2-.mu.m square sections, a standard deviation in measurements for 100 sections is 4.0% or more; and (5) the concentration of Mn in a ferrite phase is 0.90 times or less the concentration of Mn in the steel sheet, when analyzed with an electron probe microanalyzer.
2. The high strength cold rolled steel sheet according to claim 1, wherein the volume ratio of retained austenite relative to the entire structure is 5% or more, when measured by an X-ray diffraction method.
3. The high strength cold rolled steel sheet according to claim 1, wherein the hard phase comprises a mixed structure of fresh martensite and retained austenite, and at least one structure selected from the group consisting of bainitic ferrite, bainite, and tempered martensite.
4. The high strength cold rolled steel sheet according to claim 1, wherein the component composition further comprises one or more of (A) to (E) below, in percent by mass, as another element: (A) at least one selected from the group consisting of Cr: more than 0% to 1% or less and Mo: more than 0% to 1% or less; (B) at least one selected from the group consisting of Ti: more than 0% to 0,15% or less, Nb: more than 0% to 0.15% or less, and V: more than 0% to 0.15% or less; (C) at least one selected from the group consisting of Cu: more than 0% to 1% or less and Ni: more than 0% to 1% or less; (D) B: more than 0% to 0.005% or less; and (E) at least one selected from the group consisting of Ca: more than 0% to 0.01% or less, Mg: more than 0% to 0.01% or less, and REM: more than 0% to 0,01% or less.
5. A high strength electrogalvanized steel sheet in which an electrogalvanized layer is formed on the surface of the high strength cold rolled steel sheet according to claim 1.
6. A high strength galvanized steel sheet in which a galvanized layer is formed on the surface of the high strength cold rolled steel sheet according to claim 1.
7. A high strength galvannealed steel sheet in which a galvannealed layer is formed on the surface of the high strength cold rolled steel sheet according to claim 1.
8. A method for producing a high strength cold rolled steel sheet having a tensile strength of 980 MPa or higher and excellent ductility and bendability, the steel sheet being the high strength cold rolled steel sheet according to claim 1, comprising: in hot rolling a steel sheet having the component composition, coiling at a coiling temperature of 500.degree. C. to 800.degree. C., then holding for 3 hours or more at 500.degree. C. to 800.degree. C., then cooling to room temperature, soaking and holding in a temperature range from (Ac.sub.1 point+20.degree. C.) to less than Ac.sub.3 point after cold rolling, then cooling to a temperature range of 500.degree. C. or less at an average cooling rate of 10.degree. C./s or more to 500.degree. C. and at an average cooling rate of 10.degree. C./s or more at 500.degree. C. or less, then reheating to a temperature range of 250.degree. C. to 500.degree. C., holding for 30 seconds or more, and cooling to room temperature.
9. A method for producing a high strength electrogalvanized steel sheet, wherein the high strength cold rolled steel sheet obtained by the producing method according to claim 8 is fluffier subjected to electrogalvanizing.
10. A method for producing a high strength galvanized steel sheet having a tensile strength of 980 MPa or higher and excellent ductility and bendability, the steel sheet being the high strength galvanized steel sheet according to claim 6. comprising: in hot rolling a steel sheet having the component composition, coiling at a coiling temperature of 500.degree. C. to 800.degree. C., then holding for 3 hours or more at 500.degree. C. to 800.degree. C., then cooling to room temperature, soaking and holding in a temperature range from (Ac.sub.1 point+20.degree. C.) to less than Ac.sub.3 point after cold rolling, then cooling to a temperature range of 500.degree. C. or less at an average cooling rate of 10.degree. C./s or more to 500.degree. C. and at an average cooling rate of 10.degree. C./s or more at 500.degree. C. or less, then reheating to a temperature range of 250.degree. C. to 500.degree. C., holding for 30 seconds or more, galvanizing within the holding period, and then cooling to room temperature.
11. The method for producing a high strength galvanized steel sheet according to claim 10, wherein alloying in a temperature range from 450.degree. C. to 550.degree. C. is performed after the galvanization.
12. A steel sheet comprising percent by mass: C: 0.10% to 0.30%, Si: 1.2% to 3%; Mn: 0.5% to 3.0%; P: more than 0% to 0.1% or less; S: more than 0% to 0.05% or less; Al: 0.005% to 0.2%; N: more than 0% to 0.01% or less; and O: more than 0% to 0.01% or less, the balance being iron and inevitable impurities; wherein structures at a position of 1/4 sheet thickness in the steel sheet satisfy (1) to (5) below: (1) an area ratio of ferrite relative to the entire structure is from 5% or more to less than 50%, and the balance is a hard phase, when observed under a scanning electron microscope; (2) an area ratio of a mixed structure of fresh martensite and retained austenite relative to the entire structure is more than 0% to 30% or less, when observed under an optical microscope after LePera etching; (3) a region in which the concentration of Mn is enriched to 1.2 times or more the concentration of Mn in the steel sheet is present at 5 area% or more, when analyzed with an electron probe microanalyzer; (4) when the fraction of a region in which the concentration of Mn is enriched to 1.2 times or more the concentration of Mn in the steel sheet is measured in 2-m square sections, a standard deviation in measurements for 100 sections is 4.0% or more; and (5) a concentration of Mn in a ferrite phase is 0.90 times or less the concentration of Mn in the steel sheet, when analyzed with an electron probe microanalyzer.
 The present invention relates to a high strength cold rolled steel sheet and a high strength galvanized steel sheet having excellent ductility and bendability and a method for producing same. More specifically, the present invention relates to a high strength cold rolled steel sheet, a high strength electrogalvanized steel sheet, a high strength galvanized steel sheet, and a high strength galvannealed steel sheet having excellent ductility and bendability in a tensile strength range of 980 MPa or higher, and to a method for producing same suitable for efficient production of such steel sheets.
 In order to improve fuel efficiency of automobiles and transportation machines and the like, the weight of the automobiles and transportation machines needs to be reduced. For example, the weight can be effectively reduced by using high-strength steel sheets and reducing the sheet thickness. However, where the strength of steel sheets is increased, the ductility is decreased, and therefore processability is degraded. Accordingly, excellent bendability and ductility necessary for press molding are required for high-strength steel sheets. Further, to ensure etching resistance, steel sheets plated with zinc by electrogalvanizing (EG), galvanizing (GI, Galvanized Iron), or gall/annealing (GA) are often used for steel automotive parts. The electrogalvanized steel sheet, galvanized steel sheet, and galvannealed steel sheet will be sometimes referred to hereinbelow as "galvanized steel sheets". Properties required for the galvanized steel sheets are the same as those required for high-strength steel sheets.
 Examples of ultrahigh-strength steel sheets having excellent bendability which are suggested in Patent Literature 1 to 4 illustrate techniques for improving the processability of high-strength steel sheets.
 However, a steel sheet having excellent ductility and bendability in a tensile strength range of 980 MPa or higher has not yet been suggested.
 The present invention has been created with the foregoing in view,and it is an objective thereof to provide a high-strength steel sheet and a high strength galvanized steel sheet having a tensile strength of 980 MPa or higher and excellent ductility and bendability and also a method suitable for production same with good productivity, and more specifically to provide a high strength cold rolled steel sheet, a high strength electrogalvanized steel sheet, a high strength galvanized steel sheet, and a high strength galvannealed steel sheet having excellent ductility and bendability in a tensile strength range of 980 MPa or higher and a method for manufacturing same suitable for efficient production of such steel sheets.
 The gist of the present invention which can resolve the problems is that the component composition of a steel sheet comprises, in percent by mass: C: 0.10% to 0.30%; Si: 1.2% to 3%; Mn: 0,5% to 3.0%; P: more than 0% to 0.1% or less; S: more than 0% to 0.05% or less; Al: 0.005% to 0.2%; N: more than 0% to 0.01% or less; and O: more than 0% to 0.01% or less, the balance being iron and inevitable impurities, and structures at a position of 1/4 sheet thickness in the steel sheet satisfy all of (1) to (5) below. Hereinbelow, "%" in relation to the component composition of the steel sheets means "% by mass".
 (1) The area ratio of ferrite relative to the entire structure is from 5% or more to less than 50% and the balance is a hard phase, when observed under a scanning electron microscope.
 (2) The area ratio of a mixed structure of fresh martensite and retained austenite relative to the entire structure is more than 0% to 30% or less, when observed under an optical microscope after LePera etching.
 (5) The concentration of Mn in a ferrite phase is 0.90 times or less the concentration of Mn in the steel sheet, when analyzed with an electron probe microanalyzer.
 In the preferred embodiment of the present invention, the volume ratio of retained austenite relative to the entire structure is 5.sup.0 or more, when measured with an X-ray diffraction method.
 Further, it is preferred that the hard phase be composed of: the mixed structure of fresh martensite and retained austenite; and at least one structure selected from the group consisting of bainitic ferrite, bainite, and tempered martensi.
 When the present invention is implemented, in the preferred embodiment, at least any one of the following other elements is further comprised: (A) at least one selected from the group consisting of Cr: more than 0% to 1% or less and Mo: more than 0% to 1% or less; (B) at least one selected from the group consisting of Ti: more than 0% to 0.15% or less, Nb: more than 0% to 0.15% or less; and V: more than 0% to 0.15% or less; (C) at least one selected from the group consisting of Cu: more than 0% to 1% or less and Ni: more than 0% to 1% or less; (D) B: more than 0% to 0.005% or less; and (E) at least one selected from the group consisting of Ca: more than 0% to 0.01% or less, Mg: more than 0% to 0.01% or less, and REM (Rare Earth Metal): more than 0% to 0.01% or less.
 The present invention is also inclusive of a high strength electrogalvanized steel sheet in which an electrogalvanized layer is formed on the surface of the high strength cold rolled steel sheet; a high strength gair lized steel sheet in which a galvanized layer is formed on the surface of the high strength cold rolled steel sheet; and a high strength gal vannealed steel sheet in which a galvannealed layer is formed on the s face of the high strength cold rolled steel sheet.
 The method for producing the high strength cold rolled steel sheet in accordance with the present invention which can resolve the problems includes: in a step of hot rolling a steel sheet having the component composition, coiling at a coiling temperature of 500.degree. C. to 800.degree. C., then holding for 3 hours or more at 500.degree. C. to 800.degree. C., then cooling to room temperature, soak ng d holding in a temperature range from (Ac.sub.1 point+20.degree. C.) to less than Ac.sub.3 point after cold rolling, then cooling to a temperature range of 500.degree. C. or less at an average cooling rate of 10.degree. C./s or more to 500.degree. C. and at an average cooling rate of 10.degree. C./s or more at 500.degree. C. or less, then reheating to a temperature range of 250.degree. C. to 500.degree. C., holding for 30 seconds or more, and cooling to room temperature.
 In accordance with the present invention, the steel sheet obtained by the producing method is preferably further subjected to electrogalvanizing.
 Further, a method for producing the high strength cold rolled steel sheet in accordance with the present invention, which can resolve the problems, includes: in a step of hot rolling a steel sheet having the component composition, coiling at a coiling temperature of 500.degree. C. to 800.degree. C., then holding for 3 hours or more at 500.degree. C. to 800.degree. C., then cooling to room temperature, soaking and holding in a temperature range from (Ac.sub.1 point+20.degree. C.) to less than Ac.sub.3 point after cold rolling, then cooling to a temperature range of 500.degree. C. or less at an average cooling rate of 10.degree. C./s or more to 500.degree. C. and at an average cooling rate of 10.degree. C./s or more at 500.degree. C. or less, then reheating to a temperature range of 250.degree. C. to 500.degree. C., holding for 30 seconds or more, galvanizing within the holding period, and then cooling to room temperature.
 In accordance with the present invention, alloying in a temperature range from 450.degree. C. to 550.degree. C. is preferably performed after the galvanization.
 The present invention can provide a high strength cold rolled steel sheet and a high strength galvanized steel sheet having excellent ductility and bendability even at 980 MPa or more, and more particularly a high strength cold rolled steel sheet, a high strength electrogalvanized steel sheet, a high strength galvanized steel sheet, and a high strength galvannealed steel sheet that excel in the properties. With the method producing same in accordance with the present Tention, these steel sheets can be produced with good efficiency. Therefore, the high strength cold rolled steel sheet, etc., in accordance with the present invention is particularly very useful in industrial fields, such as automotive field.
 FIG. 1 is a conceptual diagram illustrating an example of a heat treatment pattern in the producing method of the present invention.
 The inventors have conducted a comprehensive study to improve ductility and bendability of high strength cold rolled steel sheets and high strength galvanized steel sheets, in particular, having a tensile strength of 980 MPa or higher.
 The results have demonstrated that ductility and bendability can be improved, while ensuring a high strength of 980 MPa or higher by optimizing a ferrite phase and a hard phase in the metal structure of the steel sheet and adequately controlling the segregation of Mn, provided that the component composition is adequately controlled.
 The reason for specifying the metal structures in the present invention is described below. A fraction measured by microscopic observations means a fraction in 100% of the entire structure of the steel sheet. As for the metal structure in the present invention, the measurement method differs depending on the metal structure. As a consequence, when the metal structures specified by the present invention are all added up, 100% is sometimes exceeded. This is because the retained y constituting the mixed structure of fresh martensite and retained austenite is measured not only by the optical microscope observations, but also by X-ray diffraction. The retained austenite can be referred to hereinbelow as "retained .gamma.", and the mixed structure of fresh martensite and retained austenite can be referred to as "MA (Martensite-Austenite Constituent) structure".
 Ferrite is a structure that increases ductility and bendability of the steel sheet. In the present invention the ductility and bendability in a high-strength region with a tensile strength of 980 MPa or higher can be increased by increasing the area fraction of ferrite. For such as effect to be demonstrated, the area ratio of ferrite is 5% or more, preferably 7% or more, more preferably 10% or more. However, where the amount of ferrite is excessively high, the strength of the steel sheet decreases and a high strength of 980 MPa or higher is difficult to ensure. Therefore, the area ratio of ferrite is set to less than 50%, preferably 45% or less, more preferably 40% or less. The area ratio of ferrite is measured by scanning electron microscope (SEM) observations at a position of 1/4 sheet thickness in the steel sheet.
 The hard phase is a structure necessary for increasing the tensile strength. In the present invention,a high strength of 980 MPa or higher can be achieved, while ensuring the presence of soft ferrite within the area ratio range, by increasing the area fraction of the hard phase. For this effect to be demonstrated, the remaining metal structure other than ferrite needs to be a hard phase. The hard phase, as referred to in the present invention, is a phase harder than the ferrite and is at least one selected from the group consisting of bainitic ferrite, bainite, tempered martensite, and MA structure. In the present invention, as indicated hereinbelow, at least the MA structure is included. Among the hard phases, bainitic ferrite, bainite, and tempered martensite are represented by values measured by SEM observations at a position of 1/4 sheet thickness in the steel sheet. The retained .gamma. is present between the laths of bainitic ferrite or included in the MA structure.
 Where the MA structure is present, the strength and ductility can be increased. Therefore, from the standpoint of increasing the strength-ductility balance, it is preferred that the area ratio of the MA structure be 3% or more, more preferably 4% or more. Meanwhile, where the area ratio of the MA structure becomes too large, benda.bility is degraded. Accordingly, in the present invention, the area ratio of the MA structure is set to 30% or less, preferably 20% or less, more preferably 15% or less.
 The fresh martensite constituting the MA structure, as referred to herein, is a product of martensitic transformation of non-transformed austenite in the process of cooling the steel sheet form the heating temperature to room temperature and is distinguished from the tempered martensite after the heating treatment. In the present invention, a location which appears white in optical microscope observations after LePera etching is taken as the MA structure. Since the fresh martensite and retained -y are difficult to distinguish from one another in optical microscope observations, the composite structure of fresh martensite and retained y is measured as the MA structure. The MA structure is represented by a value measured by optical microscope observations at a position of 1/4 sheet thickness in the steel sheet.
 In the present invention, a region in which the concentration of Mn is enriched is specified by using the concentration of Mn obtained by analysis using an electron probe microanalyzer (EPMA) with a beam diameter of 1 .mu.m or less in a 20 .mu.m.times.20 .mu.m range in the cross section of the steel sheet. The "concentration of Mn in the steel sheet", as referred to herein, is the concentration of Mn obtained by chemical analysis of the base steel sheet by an inductively coupled plasma emission spectroscopy. Therefore, the region in which the concentration of Mn is enriched to 1.2 times or more the concentration of Mn in the st sheet, as referred to herein, is a region in which the measured value of the concentration of Mn obtained by the EPMA analysis is 1.2 times or more higher than the concentration of Mn in the base steel sheet, the measurements being performed in a 20 .mu.m.times.20 .mu.m range.
 The "2-.mu.m square", as referred to herein, is a quadrangular section with a 2-.mu.m side. In the present invention, the 2 .mu.m.times.20 .mu.m EPMA measurement region is divided into 100 sections, one section being a quadrangle with a 2-.mu.m side, which are obtained by drawing vertical and horizontal lines at 2-.mu.m intervals, the area fraction of a region in which the concentration of Mn is 1.2 times or more higher is measured in each section, and the standard. deviation is determined, statistically in 100 sections.
 In the present invention, it has been found that bendability is greatly increased when the region in which the concentration of Mn is enriched to 1.2 times or more the concentration of Mn in the steel sheet is present at 5 area% or more, and the standard deviation obtained when the fraction of the region in which the concentration of Mn is enriched to 1.2 times or more in 2-.mu.m squares is 4.0% or more. The region in which the concentration of Mn in the steel sheet is enriched to 1.2 times or more can be referred to hereinbelow as the "region with the concentration of Mn of 1.2 times or more", and the standard deviation obtained when measuring the fraction of the region in which the concentration of Mn is enriched to 1.2 times or more in 2-.mu.m squares can be referred to hereinbelow as the "standard deviation of the regions in which the concentration of Mn is enriched to 1.2 times or more" or simply the "standard deviation".
 Thus, the region with the concentration of Mn of 1.2 times or more is mainly a hard phase. As the area ratio of the region with the concentration of Mn of 1.2 times or more increases, the concentration of Mn in the ferrite phase becomes relatively smaller, the hardness of the ferrite phase can be decreased, and the bendability can be increased. The larger is the segregation of Mn, the greater is the standard deviation, but the concentration of Mn in the ferrite phase, which contributes to the increase in bendability, decreases and the hardness of the ferrite phase can be reduced.
 In order to obtain such an effect, the area ratio of the region with the concentration of Mn of 1.2 times or more is set to 5.0 area % or more, preferably 5.2 area % or more, more preferably 5.5 area % or more. Meanwhile, where the ratio occupied by the region with the concentration of Mn of 1.2 times or more is too large, the Ms point of austenite decreases and the amount of the MA structure can increase. Therefore, the area ratio of the region with the concentration of Mn of 1.2 times or more is set preferably to 20 area % or less, more preferably 15 area % or less.
 Further, the standard deviation of the regions in which the concentration of Mn is enriched to 1.2 times or more is set to 4.0% or more, preferably 4.5% or more, more preferably 5.0% or more. Where the standard deviation is less than 4.0%, the distribution of Mn becomes insufficiently uniform and the decrease in hardness of the ferrite phase which contributes to the increase in bendability is insufficient. Meanwhile, no specific upper limit is set for the standard deviation, and the standard deviation is preferably 10% or less.
 Where the region with the concentration of Mn of 1.2 times or more is 5 area% or more and the standard deviation is 4.0% or more, a spheroidized hard phase can be dispersed in the ferrite phase, and the effect of increasing the strength of the steel material can be combined with the effect of increasing the bendability by the ferrite. The spheroidized hard phase is referred to hereinbelow as "globular hard phase". Here, the globular hard phase is part of the hard phase and is configured, similarly to the hard phase, of bainitic ferrite, bainite, tempered martensite, MA structure, or the like. It has been conventionally assumed that where a difference in hardness between the hard phase and ferrite phase is large, interface cracking occurs during bending and bendability is degraded, but it has been found that the interface cracking can be suppressed by the globular hard phase in the ferrite phase. In order to obtain such an effect, a smaller amount of the globular hard phase in the ferrite phase is preferred. The aspect ratio of the globular hard phase is preferably set to 3 or less, more preferably 2.5 or less, even more preferably 2 or less, and the equivalent circle diameter of the globular phase is 2 .mu.m or less, preferably 1.8 .mu.m or less, even more preferably 1.5 .mu.m or less. For the effect to be demonstrated, it is preferred that the globular hard phase take 0.70 vol % or more, more preferably 0.75 vol % or more, even more preferably 0.80 vol % or more relative to the hard phase.
 The concentration of Mn in ferrite is 0.90 times or less the concentration of Mn in the steel sheet.
 Where the concentration of Mn in the ferrite phase is too high, the hardness of the ferrite phase is not sufficiently reduced and the bendability is degraded. Therefore, the concentration of Mn in the ferrite needs to he made lower than the concentration of Mn in the steel sheet. As a consequence, the concentration of Mn in the ferrite phase is set to 0.90 times or less, preferably 0.85 times or less, more preferably 0.80 times or less the concentration of Mn in the steel sheet. Meanwhile, where the concentration of Mn in the ferrite is made too low, the ferrite hardness decreases and strength can be insufficient. For this reason, the concentration of Mn in the ferrite phase is preferably set to 0.3 times or more, more preferably to 0.4 times or more the concentration of Mn in the steel sheet. The concentration of Mn in the ferrite phase can be measured by EPMA.
 When the steel sheet is processed, the retained y receives strains, deforms, and undergoes martensitic transformation, thereby making it possible to ensure good ductility and demonstrating an effect of enhancing hardening in the deformation zone during processing and suppressing the concentration of strains. Therefore, the retained austenite is a structure necessary for improving the strength-ductility balance of the steel sheet. For such effects to be advantageously demonstrated, the volume ratio of the retained .gamma. is set preferably to 5% or more, more preferably 6% or more, even more preferably 7% or more. No specific upper limit is set for the volume ratio of the retained .gamma., but within the ranges of the component composition and producing conditions of the present invention, the volume ratio is, at the largest, 20% or less. The ratio of the retained y is a measured value determined by the X-ray diffraction at a position of 1/4 plate thickness in the steel sheet.
 The retained .gamma. is present between the laths of bainitic ferrite and included in the MA structure. Since the effect of the retained .gamma. is demonstrated regardless of the form of presence thereof, in the present invention, the retained .gamma. that can be verified during the measurements is taken as the retained y regardless of the form of presence.
 The component composition of the high-strength steel sheet in accordance with the present invention will be explained hereinbelow.
 C is an element required to ensure the strength of the steel sheet and increase the stability of the retained .gamma.. In order to ensure the tensile strength of 980 MPa or higher, the amount of C is set to 0.10% or more, preferably 0.12% or more, more preferably 0.15% or more. However, where the amount of C is too high, the strength after hot rolling rises, cracks occur during cold rolling, and weldability of the final product decreases, For this reason, the amount of C is set to 0.30% or less, preferably 0.26% or less, more preferably 0.23% or less.
 Si is a solid solution strengthening element and contributes to the strengthening of steel. Silicon also effectively suppresses the generation of carbides, effectively induces the generation of the retained .gamma., and ensures excellent TS.times.EL balance. For those effects to be advantageously demonstrated, the amount of Si is set to 1.2% or more, preferably 1.35% or more, more preferably 1,5% or more. However, where the amount of Si is too high, scale is intensively formed during the hot rolling, scale traces are transferred to the steel sheet surface, and the surface state is degraded. Pickling ability is also degraded. For this reason, the amount of Si is set to 3% or less, preferably 2.8% or less, more preferably 2.6% or less.
 Mn is an element that improves hardenability and contributes to strengthening of steel sheets. Manganese is also effective in stabilizing .gamma. and inducing the generation of retained .gamma.. For these effects to be advantageously demonstrated, the amount of Mn is set to 0.5% or more, preferably 0.6% or more, more preferably 1.0% or more, even more preferably 1.5% or more, and still more preferably to 2.0% or more. However, where the amount of Mn is too high, the strength after hot rolling increases, cracks appear during the cold rolling, and weldability of the final product is degraded. Further, where Mn is added in excess, Mn segregates and processability is degraded. For this reason, the amount of Mn is set to 3.0% or less, preferably 2.8% or less, more preferably 2.6% or less.
 P is an element that is contained unavoidably and degrades the weldability of steel sheets. For this reason, the amount of P is set to 0.1% or less, preferably 0.08% or less, more preferably 0.05% or less. Since it is desirable that the amount of P be made as small as possible, the lower limit thereof is not particularly specified, but the industrially attainable lower limit is 0.0005%.
 Similarly to P. S is also an element that is contained unavoidably and degrades the weldability of steel sheets. S also generates sulfur-containing inclusions in steel sheets and degrades the processability of steel sheet. For this reason, the amount of S is set to 0.05% or less, preferably 0.01% or less, more preferably 0.005% or less. Since it is desirable that the amount of S be made as small as possible, the lower limit thereof is not particularly specified, but the industrially attainable lower limit is 0.0001%.
 Al is an element acting as a deoxidizing agent. For the effect thereof to be advantageously demonstrated, the amount of Al is set to 0.005% or more, preferably 0.01% or more. However, where the amount of Al is too high, the weld.ability of steel sheets is greatly degraded. For this reason, the amount of Al is set to 0.2% or less, preferably 0.15% or less, more preferably 0.10% or less.
 N is an element that is contained unavoidably, causes the precipitation of nitrides in steel sheets, and contributes to the strengthening of steel sheets. From this standpoint, it is preferred that the amount of N be 0.001% or more. However, where the amount of N is too high, a large amount of nitrides precipitate, thereby causing the degradation of elongation, stretch-flangeability .lamda., and bendability. Therefore, the amount of N is set to 0.01% or less, preferably 0.008% or less, more preferably 0.005% or less.
 O is an element that is contained unavoidably, and where the amount thereof is too high, ductility and bendability during the processing are deQraded. For this reason, the amount of O is set to 0.01% or less, preferably 0.005% or less, more preferably 0.003% or less. Since it is desirable that the amount of O be made as small as possible, the lower limit thereof is not particularly specified, but the industrially attainable lower limit is 0.0001%.
 The steel sheet in accordance with the present invention has the component composition, the balance being iron and inevitable impurities. The inevitable impurities include the P, S, N, and O which can be included in steel for the reasons associated with starting materials, resources, production facilities, and the like, and also tramp elements such as Pb, Bi, Sb, and Sn. The following elements can he also actively includes as other elements within ranges in which the effects of the present invention are not adversely affected.
 Cr and Mo are elements that effectively increase quenching ability and increase the strength of steel sheets. These elements can be used individually or together. For such effects to be advantageously demonstrated,the contents of Cr and Mo are set preferably to 0.1% or more, more preferably 0.3% or more, respectivelyr. However, where the account of these elements is too high, the processability is degraded and cost is increased. For this reason, when Cr and Mo are used individually, the amount thereof is set preferably to 1% or less, more preferably 0.8% or less, even more preferably 0.5% or less. When Cr and Mo are used together, each is included within a range with the upper limit, and the total amount thereof is preferably 1.5% or less.
 Ti, Nb, and V are elements that effectively form carbide or nitride precipitates in steel sheets, increase the strength of steel sheets, and refine old y grains. These elements can be used individually or together. For these effects to be advantageously demonstrated, the amount of each of Ti, Nb, and V is set preferably to 0.005% or more, more preferably 0.010% or more. However where the amount of the elements is too high, carbides precipitate on grain boundaries and stretch-flangeability and bendability of steel sheets are degraded. Therefore, the amount of each of Ti, Nb, and V is set preferably to 0.15% or less, more preferably 0.12% or less, even more preferably 0.10% or less.
 Cu and Ni are elements effective in generating and stabilizing retained austenite and also effective in increasing the etching resistance. These elements can be used individually or together. For these effects to be advantageously demonstrated, the amount of each of Cu and Ni is set preferably to 0.05% or more, more preferably 0.10% or more. However, where Cu is comprised in excess, hot processability is degraded. For this reason, where Cu is added individually, the amount thereof is set preferably to 1% or less, more preferably 0.8% or less, more preferably 0.5% or less, Where Ni is comprised in excess, cost is increased. For this reason, the amount of Ni is set preferably to 1% or less, more preferably 0.8% or less, even more preferably 0.5% or less. Where Cu and Ni are comprised together, the effects are easily demonstrated, Furthermore, the addition of Ni suppresses the degradation of hot processability caused by the addition of Cu. Therefore, when Cu and Ni are used together, the combined amount thereof is set preferably to 1.5% or less, more preferably 1.0% or less.
 B is an element increasing quenching ability and effective in causing austenite to be stably present at room temperature. For these effects to be advantageously demonstrated, the amount of B is set preferably to 0.0005% or more, more preferably 0.0010% or more, even more preferably 0.0015% or more. However, where boron is comprised in excess, borides are generated and ductility is degraded. For this reason, the amount of B is set preferably to 0.005% or less, more preferably 0.004% or less, even more preferably 0.0035% or less.
 Ca, Mg, and REM are elements effectively refinina and dispersing inclusions in steel sheets. These elements may be used individually or in appropriate combinations of two or more thereof. For the effects thereof to be advantageously demonstrated, the amount of each of Ca, Mg and REM is set preferably to 0.0005% or more preferably to 0.0010% or more. However, where these elements are comprised in excess, castability, hot processability, or the like are degraded. For this reason, the amount of eaCh of Ca, Mg and REM is set preferably to 0.01% or less, more preferably 0.008% or less, even more preferably 0.007% or less.
 In the present invention, REM is an abbreviation of rare earth metals and is intended to include lanthanoids, that is 15 elements from La to Lu, and also scandium and yttrium.
 Methods for producing the high strength cold rolled steel sheet, high strength electrogalvanized stud sheet, high strength galvanized stud sheet, and high strength gal vannealed steel sheet in accordance with the present invention, in particular methods for producing the high strength cold rolled steel sheet and high strength galvanized steel sheet will be explained hereinbelow. In the methods for producing the high strength cold rolled steel sheet and high strength galvanized steel sheet, "in a step of hot rolling a steel sheet having the component composition, the steps of coiling at a coiling temperature of 500.degree. C. to 800.degree. C., then holding for 3 h or more at 500.degree. C. to 800.degree. C., then cooling to room temperature, soak ng uud holding in a temperature range from (Ac.sub.1 point+20.degree. C.) to less than Ac.sub.3 point after the cooling, then cooling to a temperature range of 500.degree. C. or less at an average cooling rate of 10.degree. C./s or more to 500.degree. C. and at an average cooling rate of 10.degree. C./s or more at 500.degree. C. or less" are the same Therefore, the explanation of these steps refers to both methods. The step of reheating after the "cooling to a temperature range of 500.degree. C. or less" differs among the two methods. Therefore, this step is explained separately for the high strength cold rolled steel sheet and high strength galvanized steel sheet.
 In the methods for producing the high strength cold rolled steel sheet and high strength galvanized steel sheet in accordance with the present invention, the below-described annealing is performed with respect to steel sheets obtained by hot rolling and cold rolling of the steel having the component composition. In the method for producing the high strength cold rolled steel sheet, reheating is performed after the annealing. If necessary, a high strength electrogalvanized steel sheet can be obtained by appropriately combining electrogalvanization. In the method for producing the high strength galvanized steel sheet, galvanization is performed together with reheating after the annealing. If necessary, a high-strength galvannealed steel sheet can be obtained by appropriately combining annealing processing. In accordance with the present invention, a high strength cold rolled steel sheet or a high strength galvanized steel sheet and the like having the desired structure can be obtained by appropriately controlling the producing conditions.
 For example, hot rolling is performed, as depicted in FIG. 1, by the usual method by using the steel having the component composition. In the hot rolling, for example, the steel is hot rolled such that the finish rolling temperature is Ac.sub.3 point or higher, and then coiling is performed at a coiling temperature of 500.degree. C. to 800.degree. C., Then, the steel is held for 3 h or more at 500.degree. C. to 800.degree. C., then cooled to room temperature and cold rolled. The operational upper limit for cooling after the finish rolling is about 500.degree. C./s.
 After the cold rolling, the steel is annealed by soaking and holding in a two-phase temperature range from (Ac.sub.1 point+20.degree. C.) to less than Ac.sub.3 point and then cooling to a temperature range of 500.degree. C. or less at an average cooling rate of 10.degree. C./s or more to 500.degree. C. and at an average cooling rate of 10.degree. C./s or more at 500.degree. C. or less. The method for producing a high strength cold rolled steel sheet further includes a step of reheating to a temperature range of 250.degree. C. to 500.degree. C., holding for 30 s or more, and then cooling to room temperature. The method for producing a high strength galvanized steel sheet further includes a step of reheating to a temperature range of 250.degree. C. to 500.degree. C. after cooling to the temperature range of 500.degree. C. or less, holding for 30 s or more, galvanizing within the holding period, and then cooling to room temperature.
 The reasons for specifying the conditions are described in details hereinbelow.
 As a result of coiling at a coiling temperature of 500.degree. C. to 800.degree. C. after the hot rolling and then holding for 3 h or more at 500.degree. C. to 800.degree. C., the predetermined Mn concentration distribution is generated, Mn-containing carbides are precipitated, and a hard phase spheroidized in the ferrite phase by annealing after the cold rolling is obtained. To obtain such effects, the coiling temperature is set to 500.degree. C. or more, preferably 550.degree. C. or more, more preferably 600.degree. C. or more. However, where the coiling temperature is too high, a large amount of scale or grain boundary oxidation appear in the steel sheet and pickling ability is degraded. For this reason, the coiling temperature is set to 800.degree. C. or less, preferably 750.degree. C. or less, more preferably 700.degree. C. or less. Further, the temperature range of holding after the coiling is set to 500.degree. C. or more, preferably 510.degree. C. or more, more preferably 520.degree. C. or more, even more preferably 550.degree. C. or more, and still more preferably 580.degree. C. or more. Meanwhile, where the holding temperature is too high, a large amount of scale, grain boundary oxidation, or the like, can appear in the steel sheet and pickling ability can be degraded in the same manner as when the coiling temperature is too high. For this reason, the holding temperature is set to 800.degree. C. or less, preferably 780.degree. C. or less, even more preferably 750.degree. C. or less, still more preferably 700.degree. C. or less. The holding time in this temperature range is 3 h or more, preferably 4 h or more, more preferably 5 h or more, even more preferably 7 h or more, and still more preferably 10 h or more Meanwhile, where the holding time is too long, a large amount of scale, grain boundary oxidation, or the like, can appear in the steel sheet and pickling ability can be degraded in the same manner as when the coiling temperature is too high. For this reason, the holding time is set preferably to 72 h or less, more preferably 60 h or less.
 In the present invention, holding at a predetermined temperature it the concept such that the holding is not necessarily continued at the same temperature and the temperature may change, provided it is within a predetermined temperature range. For example, holding may be isothermal holding within a range of the holding temperature and is inclusive of variations within this range, that is, a decrease in temperature, increase in temperature caused by heating, and increase in temperature caused by recuperation associated with transformation.
 In the present invention, cooling to room temperature is performed after holding for the preset time in the temperature range, but the cooling rate in this step is not particularly limited. For example, air cooling may be performed.
 After the hot rolling, optional pickling is performed and then cold rolling is performed at a cold rolling draft of about 30% to 80%.
 In the annealing step after the cold rolling, soaking and holding are performed in a two-phase region in a temperature range from (Ac.sub.1 point+20.degree. C.) to less than Ac.sub.3 point, and then cooling is performed to a temperature range of 500.degree. C. or less by cooling at an average cooling rate of 10.degree. C./s or more in a temperature range to 500.degree. C. and then cooling at an average cooling rate of 10.degree. C./s or more in a temperature range from 500.degree. C. or less.
 As a result of soaking and holding in a two-phase region in a temperature range from (Ac.sub.1 point+20.degree. C.) to less than Ac.sub.3 point, it is possible to ensure the desirable amount of ferrite, while maintaining, the Mn concentration distribution in accordance with the present invention. Where the holding and soaking temperature is lower than Ac.sub.1 point+20.degree. C., the amount of ferrite in the metal structure of the finally obtained steel sheet is too large and sufficient strength cannot be ensured. For this reason, the holding and soaking temperature is set equal to or higher than Ac.sub.1 point+20.degree. C., preferably equal to or higher than Ac.sub.1 point+25.degree. C., more preferably equal to or higher than Ac.sub.1 point+50.degree. C., and even more preferably equal to or higher than Ac.sub.1 point+80.degree. C. Meanwhile where the Ac.sub.3 point is reached or exceeded, ferrite cannot be sufficiently generated and grown during soaking and holding, the ductility decreases, the Mn concentration distribution becomes even, and the amount of the globular hard phase generated in the ferrite phase decreases. For this reason the holding and soaking temperature is set less than Ac.sub.3 point, preferably equal to or less than Ac.sub.3 point-5.degree. C., more preferably equal to or less than Ac.sub.3 point-10.degree. C., and even more preferably equal to or less than Ac.sub.3 point-20.degree. C.
 The average temperature increase rate when the temperature rises in the soaking and holding temperature range is not particularly limited and can be selected as appropriate. For example, the average temperature increase rate may be about 0.5.degree. C./s to 50.degree. C./s.
 In the present invention, the holding time in the soaking and holding temperature range is not particularly limited. However, where the holding time is too short, the processed structure remains and the ductility of steel can decrease. For this reason, the holding time is set preferably to 40 s or more, more preferably to 60 s or more. Meanwhile, where the holding time is too long, the enrichment of the austenite phase with Mn is advanced, the Ms point decreases, and the amount of MA structure can be increased. For this reason, the holding time is set preferably to 3600 s or less, more preferably 3000 s or less.
 As mentioned hereinabove, in the present invention, holding at a predetermined temperature it the concept such that the holding is not necessarily continued at the same temperature and the temperature may change, provided it is within a predetermined temperature range. For example, when holding is performed at the soaking and holding temperature, it may be isothermal holding within a range from a temperature equal to or higher than Ac.sub.1+20.degree. C. to less than Ac.sub.3, and the holding temperature may vary within this range.
 The Ac.sub.1 point and Ac.sub.3 points can be calculated from formulas (a) and (b) below which are described in in "The Physical Metallurgy of Steels", published by Maruzen on May 31, 1985, p. 273. In the formulas, the square parentheses represent the amount of each element in mass %. The amount of elements which are not contained in the steel sheet may be calculated as 0 mass %.
 After the soaking and holding, cooling is performed in a temperature range to 500.degree. C. at an average cooling rate of 10.degree. C./s or more. By controlling the cooling rate from the soaking and cooling temperature, it is possible to suppress the generation of ferrite with a high concentration of Mn and reduce the amount of generated ferrite. Where the average cooling rate is low, ferrite with a high concentration of Mn is generated during the cooling, the bendability can be degraded, and the strength can be reduced. For this reason, the average cooling rate is set to 10.degree. C./s or more, preferably 15.degree. C./s or more, more preferably 20.degree. C./s or more. No upper limit is set for the average cooling rate, and cooling may be performed with water or oil.
 After the cooling at the average cooling rate in a temperature range to 500.degree. C., cooling is performed at 500.degree. C. or less at an average cooling rate of 10.degree. C./s or more. By setting the average cooling rate at 500.degree. C. or less to 10.degree. C./s or more, it is possible to suppress the generation of soft high-temperature bainite, suppress the auto-tempering martensite, and increase the strength. For such effects to be obtained, the average cooling rate at 500.degree. C. or less is set to 10.degree. C./s or more, preferably 15.degree. C./s or more, more preferably 20.degree. C./s or more. No upper limit is set for the average cooling rate, and cooling may be performed with water or oil.
 The average cooling rate to 500.degree. C. and the average cooling rate at 500.degree. C. or less may be the same or different and may be adjusted, as appropriate, within the ranges.
 The cooling stop temperature when the cooling is performed at an average cooling rate of 10.degree. C./s or more at 500.degree. C. or less is in a temperature range of 500.degree. C. or less. Where the cooling stop temperature is higher than 500.degree. C., the amount of the hard phase decreases, the strength cannot be ensured, and the amount of the MA structure increases, thereby degrading the bendability. For this reason, the cooling stop temperature is set to 500.degree. C. or less, preferably 400.degree. C. or less, more preferably 350.degree. C. or less, even more preferably 300.degree. C. or less. No lower limit is set for the cooling stop temperature, and it can be as low as room temperature to simplify the operations.
 The reheating step will be explained hereinbelow separately for the method for producing a cold-rolled steel sheet and the method for producing a galvanized steel sheet.
 After the cooling in a temperature range of 500.degree. C. or less has been stopped, reheating is performed to a temperature range of 250.degree. C. to 500.degree. C., followed by holding for 30 s or more, and cooling to room temperature.
 By reheating to a temperature range of 250.degree. C. to 500.degree. C. and holding for 30 s or more after the cooling stop, it is possible to temper the hard phase such as martensite and cause the transformation of the non-transformed austenite. When the reheating is not performed or the holding temperature is too low, the tempering of the hard phase does not advance, high-density dislocations are generated, a large amount of the MA structure remains, and the bendability is degraded. For this reason, the reheating temperature is set to 250.degree. C. or more, preferably 300.degree. C. or more, more preferably 350.degree. C. or more. Meanwhile, where the holding temperature is too high, the strength decreases. For this reason, the reheating temperature is set to 500.degree. C. or less, preferably 470.degree. C. or less, more preferably 450.degree. C. or less. In the present invention, the "reheating" is literary the heating, that is, increase in temperature, from the cooling stop temperature to 500.degree. C. or less. Therefore, the reheating temperature is a temperature higher than the cooling stop temperature, and even in a temperature range of 250.degree. C. to 500.degree. C., the isothermal holding in which the cooling stop temperature is equal to the reheating temperature, or the process of cooling from the cooling stop temperature to an even lower temperature is not included in the reheating.
 Further, where the holding time in the reheating temperature range is too short, the hard phase cannot be sufficiently tempered, and the non-transformed austenite cannot he transformed. For this reason, the holding time is set to 30 s or more, preferably 50 s or more, more preferably 100 s or more, even more preferably 200 s or more. Meanwhile, no upper limit is set for the holding time, but where the holding time is too long, the productivity decreases. In addition, the strength also decreases. For this reason, the holding time is set preferably to 1500 s or less, more preferably 1000 s or less.
 After holding for the predetermined time in the reheating temperature range, cooling is performed to room temperature. The average cooling rate in this case is not particularly limited, and the cooling may he performed at an average cooling rate of preferably 0.1.degree. C./s or more, more preferably 0.4.degree. C./s or more, and preferably 200.degree. C./s or less, more preferably 150.degree. C./s or less.
 In accordance with the present invention, an electrogalvanized layer may be formed on the surface of the obtained steel sheet.
 A method for forming the electrogalvanized layer is not particularly limited, and the usual electrogalvanizing method can be used. For example, when an electrogalvanized steel sheet is produced, the electrogalvanizing treatment can be performed by immersing in a zinc solution at 55.degree. C. and allowing an electric current to flow in the solution. The amount plated on one side is not particularly limited and can be about 10 g/m.sup.2 to 100 g/m.sup.2 when an electrogalvanized steel sheet is produced.
 After the cooling in a temperature range of 500.degree. C. or less has been stopped, reheating is performed to a temperature range of 250.degree. C. to 500.degree. C., followed by holding for 30 s or more, galvanizing in the holding time, and then cooling to room temperature.
 By reheating to a temperature range of 250.degree. C. to 500.degree. C. and after the cooling stop and holding for 30 s or more, it is possible to temper the hard phase such as martensite and cause the transformation of the non-transformed austenite. When the reheating is not performed or the holding temperature is too low, the tempering of the hard phase does not advance, high-density dislocations are generated, a large amount of the MA structure remains, and the bendability is degraded. For this reason, the reheating temperature is set to 250.degree. C. or more, preferably 300.degree. C. or more, more preferably 350.degree. C. or more. Meanwhile, where the holding temperature is too high, the strength decreases. For this reason, the reheating temperature is set to 500.degree. C. or less, preferably 470.degree. C. or less, more preferably 450.degree. C. or less. In the present invention, the "reheating" is literary the heating, that is, increase in temperature, from the cooling stop temperature to 500.degree. C. or less. Therefore, the reheating temperature is a temperature higher than the cooling stop temperature, and even in a temperature range of range of 250.degree. C. to 500.degree. C., the isothermal holding in which the cooling stop temperature is equal to the reheating temperature, or the process of cooling from the cooling stop temperature to an even lower temperature is not included in the reheating.
 Further, where the holding time in the reheating temperature range is too short, the hard phase cannot be sufficiently tempered, and the non-transformed austenite cannot be transformed. For this reason, the holding time is set to 30 s or more, preferably 50 s or more, more preferably 100 s or more, even more preferably 200 s or more. Meanwhile, no upper limit is set for the holding time, but where the holding time is too long, the productivity decreases. In addition, the strength also decreases. For this reason, the holding time is set preferably to I50( )s or less, more preferably 1000 s or less.
 In the present invention a galvanized layer is formed on the steel sheet surface by galvanization in the holding time of 30 s or more in the reheating temperature range. In the present invention, the galvanization and holding in the reheating temperature range can be performed together. Thus, in order to control as appropriate, for example, the metal structure or strength by reheating, it is necessary that galvanization be performed within the holding time period in the reheating time range. A method for forming the galvanized layer is not particularly limited and the usual galvanization method can be used. For example, the galvanization may be performed by dipping the steel sheet in a plating bath which has been adjusted to the reheating temperature range. The plating time may be the holding time and may be adjusted, as appropriate, such as to ensure the desired amount of plated metal. For example, the plating time is preferably 1 s to 10 s.
 (iii) the treatment is performed in the order of: only heating, galvanization, only heating.
 Where the reheating temperature in the case in which only heating is performed is different from the galvanization temperature, that is, the plating bath temperature, heating or cooling from one temperature to the other temperature can be included. Examples of the heating method include furnace heating and induction heating.
 When the galvannealed layer is formed on the steel sheet surface, alloying may be performed after the galvanization. The alloying temperature is not particularly limited, but where the alloying temperature is too low, the alloying does not sufficiently advance. For this reason, the alloying temperature is preferably 450.degree. C. or more, more preferably 460.degree. C. or more, and even more preferably 480.degree. C. or more. Meanwhile, where the alloying temperature is too high, the alloying advances too much, the concentration of Fe in the plated layer increases, and the adhesion of the plated layer is degraded. For this reason, the alloying temperature is preferably 550.degree. C. or less, more preferably 540.degree. C. or less, and even more preferably 530.degree. C. or less. The time of the alloying treatment is not particularly limited and may be adjusted to obtain the desired alloying. The alloying treatment time is preferably from 10 s to 60 s. Since the alloying treatment is performed after holding for the predetermined time in the reheating temperature range, the alloying treatment time is not included in the time of holding in the reheating temperature range.
 After holding for the predetermined time in the reheating temperature range, cooling is performed to room temperature. The average cooling rate in this case is not particularly limited, and the cooling may be performed at an average cooling rate of preferably 0.1.degree. C./s or more, more preferably 0,4.degree. C./s or more, and preferably 200.degree. C./s or less, more preferably 150.degree. C./s or less.
 The technique of the present invention can be particularly advantageously applied to thin steel sheets with a thickness of 6 mm or less.
 The objects of the high strength cold rolled steel sheet and high strength galvanized steel sheet in accordance with the present invention are steel sheets with a tensile strength of 980 MPa or more, preferably 1000 MPa or more, and more preferably 1010 MPa or more. The ductility is represented by the balance of strength and ductility, that is, by "(tensile strength in MPA units).times.(ductility in % units)", and is set preferably to 15,000 MPa-% or more, more preferably to 15,100 MPa-% or more, even more preferably to 15,200 MPa-% or more. The bendability is represented by the balance of strength and VDA. (V rband der Automobilindustrie) bending angle determined by the method disclosed in the below-described examples, that is, by "(tensile strength in MPA units).times.(VDA bending angle in degree units)", and is set preferably to 100,000 MPa-.degree. or more, more preferably to 100,500 MPa-.degree. or more, even more preferably to 101,000 MPa-.degree. or more.
 This application claims the benefit of priority to Japanese Patent Application No. 2014-053399 filed on Mar. 17, 2014, Japanese Patent Application No. 2014-053400 filed on Mar. 17, 2014, and Japanese Patent Application No. 2014-192757 filed on Sep. 22, 2014. The entire contents of Japanese Patent Application No. 2014-053399 filed on Mar. 17, 2014, the contents of Japanese Patent Application No. 2014-053400 filed on Mar. 17, 2014, and the entire contents of Japanese Patent Application No. 2014-192757 filed on Sep. 22, 2014 are incorporated herein by reference.
 The present invention will be explained hereinbelow in greater detail by the examples thereof, but the present invention is not limited to the examples and can be practiced with modifications adaptable to the purposes described above and below, and all those modification are also included in the technical scope of the present invention.
 A slab was heated to 1250.degree. C. and hot rolled to a sheet thickness of 2.3 mm at a draft ratio of 90% and a finish rolling temperature of 920.degree. C. Then, the steel sheet was cooled to the "Coiling temperature (.degree. C.)" shown in Table 2 or Table 3 at an average cooling rate of 30.degree. C./s from this temperature, coiled, and then held under the conditions of the "Holding temperature 1 (.degree. C.)" and "Holding time (h)" shown in Table 2, or the "Holding start temperature (.degree. C.)", "Holding end temperature (.degree. C.)", and "Holding time (h)" shown in Table 3, A hot-rolled steel sheet was then produced by air cooling to room temperature.
 The obtained hot-rolled steel sheet was pickled to remove the surface scale, and then cold rolling was performed to produce a cold-rolled steel sheet with a thickness of1:2 mm.
 A sample steel was produced by soaking/holding.fwdarw.cooling.fwdarw.reheating the obtained cold-rolled steel sheet under the conditions shown in Table 2 and Table 3. No. 32 in Table 2 is a comparative example in which no reheating was performed. Instead of reheating the steel sheet was cooled from the cooling stop temperature of 480.degree. C. to 350.degree. C. and then held for 300 s at this temperature. This process is indicated in the reheating column.
 In the tables, the soaking and holding temperature is represented by "Soaking temperature (.degree. C.)", the average cooling rate to 500.degree. C. after the soaking is represented by "Average cooling rate 1 (.degree. C./s)", the cooling rate at 500.degree. C. or less is represented by "Average cooling rate 2 (.degree. C./s), the cooling stop temperature is represented by "Cooling stop temperature (.degree. C.)", the holding temperature during reheating after the cooling stop is represented by "Reheating holding temperature (.degree. C.)", and the holding time at the holding temperature is represented by "Reheating holding time (s)". In the present example, the holding time at the soaking and holding temperature was set to 100 s to 600 s. After the "Reheating holding time (s)" at the "Reheating holding temperature (.degree. C.)", the steel sheet was allowed to cool to room temperature obtain a sample steel. The cold-rolled steel sheets which were not subjected to the below-described electrogalvanization were assigned with a "CR" symbol in the "Product type" column in the tables.
 Electrogalvanized steel sheets were obtained from some of the sample steels by immersing in a zinc plating bath at 55.degree. C., electrogalvanizing at a current density of 30 A/dm.sup.2 to 50 A/dm.sup.2, and then washing with water and drying. The amount of zinc plated on one side was 10 g/m.sup.2 to 100 g/m.sup.2. In the plating treatment, sample steels having an electrogalvanized layer on the surface were obtained by performing washing treatment such as degreasing by immersion in an appropriate aqueous alkali solution, washing with water, and pickling. The electrogalvanized steel sheets were assigned with an "EG" symbol in the "Product type" column in the tables.
 The metal structures, concentration of Mn, and mechanical properties were evaluated, as described hereinbelow, for each sample steel. The results are shown in Table 4 and Table 5.
 The area ratio of ferrite, the area ratio of the hard phase, the area ratio of the MA structure, the volume ratio of the retained .gamma., and the ratio of the globular hard phase were measured in the following manner. Thus, the cross section of the sample steel was polished and etched as described hereinabove, and the position at 1/4 of the sheet thickness was then observed under an optical microscope or a scanning electron microscope. The ratio of each structure was measured by image analysis of the metal structure image captured with the optical microscope or SEM. The details are described below.
 The polished sample was subjected to nital etching, a total of 3 fields of view, each field of view having a size of 100 .mu.m.times.100 .mu.m, were observed under the SEM with a magnification of 1000, the area ratio of ferrite was measured with a point algorithm method with 20.times.20 grating points at a grating spacing of 5 .mu.m, and the average value for the 3 fields of view was calculated. The results are shown in "Ferrite (area %)" in the tables. The hard phase in the ferrite phase is excluded from the a ratio of ferrite.
 The structures other than the ferrite are a hard phase, and a value obtained by subtracting the ferrite area ratio from 100 area % of the observation fields of view was taken as the area ratio of the hard phase. The results are shown in "Hard phase (area %)" in the tables. The structures of the hard phase were also observed, and it was confirmed that the hard phase is at least one selected from the group including bainitic ferrite, bainite, tempered martensite, retained .gamma., and MA structure.
 The polished sample was subjected to LePera etching, a total of 3 fields of view, each field of view having a size of 100 .mu.m.times.100 .mu.m, were observed under the optical microscope with a magnification of 1000, the area ratio of the MA structure was measured with a point algorithm method with 20.times.20 grating points at a grating spacing of 5 .mu.m, and the average value for the 3 fields of view was calculated. The results are shown in "MA (area %)" in the tables. A zone which was colored white in the LePera etching was observed as the MA structure.
 Volume ratio of retained .gamma.
 The sample was polished to a position of 1/4 sheet thickness with sandpapers #1000 to #1500, and the surface was then electrolytically polished to a depth of 10 .mu.m to 20 .mu.m, followed by measurements using an X-ray diffraction device RINT1500 produced by Rigaku Corporation. More specifically, a Co target was used, the measurements were performed in a 2.theta. range of 40.degree. to 130.degree. at an output of 40 kV-200 mA, and the retained .gamma. was quantitatively measured from the obtained diffraction peaks (110), (200), (211) of the bcc (.alpha.) and diffraction peaks (111), (200), (220), and (311) of fcc (.gamma.). The results are presented in "Retained .gamma. (vol %)" in the tables.
 The sample steel was cut along the transverse cross section, embedded in a resin, and polished. The concentration of Mn was then measured in a 20 .mu.m.times.20 .mu.m range by using EPMA under the condition of a beam diameter being 1 .mu.m of less. The obtained concentration of Mn was divided by the concentration of Mn in the steel sheet determined by chemical analysis performed by inductively coupled plasma emission spectroscopy, and the ratio of the region in which the concentration of Mn was enriched to 1.2 times or more the concentration of Mn in the steel sheet was determined. Then, the region with the concentration of Mn of 1.2 times or more and the region with the concentration of Mn of less than 1.2 times were represented by different colors, and the area percentage of the region with the concentration of Mn of 1.2 times or more was determined. The results are presented in "Area ratio of region with concentration of Mn of 1.2 times (%)" in the tables.
 An image colored according to the concentration of Mn in the steel sheet was divided into 100 sections, each section being a 2-.mu.m square, the fraction of the region in which the concentration of Mn was enriched to 1.2 times or more the concentration of Mn in the steel sheet was measured in each section, and the standard deviation of 100 sections was determined. The results are presented in "Standard deviation of regions with concentration of Mn of 1.2 times (area %)" in the tables.
 The concentration of Mn was observed under a SEM in the same 20 .mu.m.times.20 .mu.m fields of view measured by the EPMA analysis. The ferrite grains and distribution of Mn concentration therein were identified by comparing the ERMA analysis results with the SEM image. The intersection point of the long and short axes of a ferrite grain was taken as the center position of the ferrite grain, and the concentration of Mn in the center position was taken as the concentration of Mn of this ferrite grain. The concentration of Mn at the ferrite grain center position in 20 .mu.m.times.20 .mu.m ranges was identified by the method, and the ratio of the concentration of Mn in the ferrite phase in the concentration of Mn in the steel sheet was determined by dividing the concentration of Mn of the ferrite grain the highest concentration of Mn among the 20 .mu.m.times.20 .mu.m ranges by the concentration of Mn in the steel sheet. In the present invention, a 20 .mu.m.times.20 .mu.m field was taken as one field of view, and the average value for 3 fields of view was taken as the ratio of the concentration of Mn in the ferrite phase in the concentration of Mn in the steel sheet. The results are presented in "Ratio of concentration of Mn in ferrite phase" in the tables.
 A tensile test was performed using a test piece No. 5 stipulated by JIS Z2201, and the tensile strength and ductility were measured. The test piece was cut out from the sample steel such that the longitudinal direction was orthogonal to the rolling direction. A strength-ductility balance was calculated from the obtained tensile strength and ductility. In the tables, the tensile strength, ductility, and strength-ductility balance are presented by "TS (MPa)", "EL (%)", and "TS.times.EL (MPa. %)".
 In the present invention, when the tensile strength was 980 MPa or higher, the strength was high and evaluated as acceptable, and when the tensile strength was less than 980 MPa, the strength was insufficient and evaluated as unacceptable.
 Further, the ductility was evaluated by the strength-ductility balance, and where TS.times.EL was 15,000 MPa-% or higher, the ductility was evaluated as excellent and acceptable. Where TS.times.EL was less than 15,000 MPa-%, the ductility was evaluated as poor and unacceptable.
 The bendability was evaluated under the following measurement conditions on the basis of the VDA standard "VDA238-100" of Verband der Deutschen Automobilindustrie (the German Association of the Automotive Industry). In the present invention, the bending angle was determined by converting the displacement under the maxim load, which was obtained by the bending test, into the angle stipulated by the VDA standard. The results are presented in "VDA bending angle (.degree.)" in the table. The bendability was evaluated from the tensile strength and bending angle. The results are presented in "TS.times.VDA (MPa-.degree.)" in the tables, Where TS.times.VDA was 100,000 MPa-.degree. or more, the bendability was evaluated as excellent and acceptable, Where TS.times.VDA was less than 100,000 MPa-.degree., the bendability was evaluated as insufficient and and acceptable.
TABLE-US-00001 TABLE 1 Steel Component composition (mass %) Ac.sub.1 Ac.sub.3 grade C Si Mn P S Al N O Cr Mo Ti Nb V Cu Ni B Ca Mg REM (.degree. C.) (.degree. C.) A 0.23 1.5 2.4 0.01 0.001 0.04 0.001 0.001 0.015 741 837 B 0.19 2.0 2.6 0.01 0.001 0.04 0.003 0.001 0.015 753 862 C 0.23 1.8 2.2 0.01 0.002 0.04 0.003 0.001 0.015 752 856 D 0.23 1.5 2.0 0.01 0.001 0.04 0.004 0.001 0.015 745 849 E 0.19 1.5 2.6 0.01 0.001 0.04 0.003 0.001 739 834 F 0.23 2.2 2.0 0.01 0.001 0.03 0.004 0.001 0.3 771 867 G 0.17 1.35 2.2 0.01 0.001 0.03 0.004 0.001 0.3 739 849 H 0.17 1.35 2.0 0.02 0.002 0.03 0.005 0.001 0.015 741 853 I 0.15 1.35 2.2 0.01 0.002 0.04 0.003 0.001 0.015 739 850 J 0.17 1.85 2.3 0.01 0.001 0.04 0.004 0.001 0.1 752 861 K 0.19 1.85 2.2 0.02 0.001 0.05 0.003 0.001 0.3 748 868 L 0.21 1.85 2.1 0.01 0.002 0.04 0.005 0.001 0.1 0.3 749 853 M 0.22 1.35 2.1 0.01 0.001 0.05 0.004 0.001 0.015 0.0025 740 845 N 0.17 1.35 2.2 0.02 0.001 0.04 0.004 0.001 0.001 739 851 O 0.22 1.5 2.4 0.02 0.001 0.06 0.004 0.001 0.015 0.001 741 854 P 0.22 1.5 2.4 0.01 0.001 0.07 0.004 0.001 0.3 0.015 0.001 746 848 Q 0.19 2.8 2.0 0.01 0.001 0.01 0.003 0.001 0.015 0.0025 783 904 R 0.26 1.2 0.8 0.01 0.002 0.04 0.003 0.001 1.0 0.3 0.1 766 856 S 0.15 2.6 2.6 0.01 0.002 0.06 0.003 0.001 0.3 0.015 0.0025 0.001 771 916 T 0.12 1.8 2.8 0.01 0.001 0.07 0.004 0.001 0.9 0.9 0.1 0.0015 0.001 761 888 U 0.35 1.2 2.6 0.02 0.001 0.05 0.005 0.001 730 800 V 0.26 1.5 3.5 0.01 0.001 0.04 0.002 0.001 729 792 W 0.12 2 2.2 0.02 0.003 0.05 0.006 0.001 758 897 X 0.19 1.9 2.2 0.01 0.001 0.05 0.004 0.001 755 867 Y 0.21 1.8 2.1 0.02 0.001 0.05 0.003 0.001 753 868 Z 0.23 1.3 2.1 0.02 0.002 0.04 0.004 0.001 738 838 AA 0.28 1.2 2.0 0.02 0.002 0.05 0.002 0.001 737 830 AB 0.23 1.8 2.4 0.02 0.002 0.05 0.004 0.001 750 855 AC 0.18 1.5 2.8 0.01 0.001 0.03 0.001 0.001 737 826 * REM in steel grade P is Y: 0.0005, Ce: 0.0005.
 Tables 1 to 5 lead to the following conclusions. The steel sheets of Tests No. 1 to 18, 20, 24 to 26, 28, 29, and 33 to 43 which were produced using the steel grades A to T and W to AC having the component composition in accordance with the present invention and under the annealing conditions stipulated by the present invention had excellent ductility and bendability in a tensile strength range of 980 MPa and higher.
 By contrast, other steel sheets did not have the component compositions or were not produced under the conditions stipulated by the present invention, as described hereinbelow, and the desired properties could not be obtained.
 The amount of C in steel grade U and the amount of Mn in steel grade V in Table 1 exceeded the respective upper limits of the present invention and fracture occurred during cold rolling. As a result, sample steels could not be produced.
 Test No. 19 is an example in which no reheating was performed after cooling to 500.degree. C. or less. The amount of MA structure increased and the bendability was degraded.
 Test No. 21 is an example in which no holding at the predetermined temperature was performed after coiling. The standard deviation of the regions with the concentration of Mn of 1.2 times was low and the bendability was degraded.
 Test No. 22 is an example in which the coiling temperature and holding temperature were low. The area ratio of the region with the concentration of Mn of 1.2 times and the standard deviation of the regions with the concentration of Mn of 1.2 times were low and the bendability was degraded.
 Test No. 23 is an example in which the time of holding at the predetermined temperature after coiling was short. The standard deviation of the regions with the concentration of Mn of 1.2 times was low and the bendability was degraded.
 Test No. 27 is an example in which the soaking temperature was high. No ferrite was generated and the area ratio of the region with the concentration of Mn of 1.2 times and. the standard deviation of the regions with the concentration of Mn of 1.2 times were low. As a result, the ductility was degraded.
 Test No. 30 is an example in which the cooling stop temperature was high. The amount of ferrite and MA structure increased, the strength was low, and the ductility and bendability were degraded.
 Test No. 31 is an example in which the cooling rate to 500.degree. C. was low. The concentration of Mn in the ferrite phase became too high. As a result, the bendability was degraded.
 The steels with the component compositions presented in Table 6 were melted and cold-rolled steel sheets were produced by hot rolling.fwdarw.cold rolling.fwdarw.continuous annealing undo the below-described conditions. In the steels with the component compositions presented in Table 6, the balance is iron and inevitable impurities. The empty column means that the element was not added.
 A slab was heated to 1250.degree. C. and hot rolled to a sheet thickness of 2,3 mm at a draft ratio of 90% and a finish temperature of 920.degree. C. Then, the steel sheet was cooled to the "Coiling temperature (.degree. C.)" shown in Table 7 or Table 8 at an average cooling rate of 30.degree. C./s from this temperature, coiled, and then held under the conditions of the "Holding temperature 1 (.degree. C.)" and "Holding time (h)" shown in Table 7, or the "Holding start temperature (.degree. C.)", "Holding end temperature (.degree. C.)", and "Holding time (h)" shown in Table 8, A hot-rolled steel sheet was then produced by air cooling to room temperature.
 The obtained hot-rolled steel sheet was pickled to remove the surface scale, and then cold rolling was performed to produce a cold-rolled steel sheet with a thickness of 1.2 mm.
 A sample steel was produced by soaking/holding.rarw.cooling.rarw.reheating.rarw.plating the obtained cold-rolled steel sheet under the conditions shown in Table 9 and Table 10. No. 29 in Table 7 is a comparative example in which no reheating was performed. Instead of reheating the steel sheet was cooled from the cooling stop temperature of 470.degree. C. to 400.degree. C. and then held for 45 s at this temperature. This process is indicated in the reheating column.
 In the tables, the soaking and holding temperature is represented by "Soaking temperature (.degree. C.), the average cooling rate to 500.degree. C. after the soaking is represented by "Average cooling rate 1 (.degree. C./s)", the cooling rate at 500.degree. C. or less is represented by "Average cooling rate 2 (.degree. C.)", the cooling stop temperature is represented by "Cooling stop temperature (.degree. C.)", the holding temperature during reheating after the cooling stop is represented by "Reheating holding temperature (.degree. C.)", the holding time at the holding temperature is represented by "Reheating holding time (s)", the temperature of the plating bath is represented by "Plating bath temperature (.degree. C.)", and the plating treatment time is represented by "Galvanization treatment time (s)". The "Reheating holding time (s)" is a total time including, the "galvanization treatment time (s)".
 After holding for about ("Reheating holding time (s)"--"Galvanization treatment time (s)") at the "Reheating holding temperature (.degree. C.)", the steel sheet is dipped into a galvanization bath and the galvanized layer is formed for the "Galvanization treatment time (s)", in Tests No. 31 and 33, the steel was immersed into the galvanization bath after heating from the "Reheating holding temperature (.degree. C.)" to "Plating bath temperature (.degree. C.)", which are shown in Table 8, immediately prior to dipping into the galvanization bath. Some steel sheets were alloyed after the galvanization treatment. In the tables, the alloying temperature is represented by "Alloying temperature (.degree. C.)" and the holding time at the alloying temperature is represented by "Alloying treatment time (s)". After holding for the predetermined time, the steel sheet was allowed to cool to room temperature to obtain a sample steel. The steel sheets which were subjected only to galvanization were assigned with a "GI" symbol and the steel sheets that were also alloyed were assigned with a "GA" symbol in the "Product type" column in the tables.
 The metal structures, concentration of Mn, and mechanical properties were evaluated, as described hereinbelow, for each sample steel in the same manner as in Example 1. The results are shown in Table 9 and Table 10.
TABLE-US-00006 TABLE 6 Steel Component composition (mass %) Ac.sub.1 Ac.sub.3 grade C Si Mn P S Al N O Cr Mo Ti Nb V Cu Ni B Ca Mg REM (.degree. C.) (.degree. C.) A 0.21 1.85 2.1 0.01 0.001 0.04 0.004 0.001 754 860 B 0.17 2 2.3 0.01 0.001 0.03 0.005 0.001 757 866 C 0.21 1.2 2 0.01 0.002 0.04 0.005 0.001 737 834 D 0.22 1.5 1.9 0.01 0.001 0.04 0.004 0.001 0.4 753 843 E 0.22 1.5 1.9 0.01 0.001 0.04 0.004 0.001 0.3 746 857 F 0.22 1.8 1.8 0.01 0.001 0.04 0.003 0.001 0.02 756 872 G 0.19 1.5 2.1 0.01 0.002 0.03 0.004 0.001 0.02 744 845 H 0.15 2.6 2.4 0.01 0.001 0.05 0.004 0.001 0.02 773 905 I 0.23 1.35 2 0.01 0.001 0.03 0.003 0.001 0.2 0.2 738 825 J 0.19 2 2 0.02 0.001 0.04 0.003 0.001 0.003 760 881 K 0.26 1.35 2.2 0.01 0.002 0.04 0.004 0.001 0.002 739 824 L 0.12 1.8 2.8 0.01 0.001 0.04 0.003 0.001 0.003 745 859 M 0.26 1.8 1.6 0.01 0.001 0.05 0.003 0.001 0.002 758 866 N 0.19 1.85 2.2 0.02 0.001 0.04 0.002 0.001 0.03 0.2 753 876 O 0.35 1.2 2.6 0.01 0.001 0.04 0.004 0.001 730 789 P 0.26 1.5 3.5 0.01 0.001 0.04 0.004 0.001 729 792 Q 0.21 2 1.8 0.01 0.001 0.01 0.006 0.001 762 863 R 0.22 1.9 2.4 0.02 0.001 0.04 0.004 0.001 753 858 S 0.19 2.4 1.8 0.02 0.001 0.03 0.004 0.001 774 901 T 0.16 1.6 1.6 0.01 0.002 0.03 0.003 0.001 743 854 U 0.19 1.95 2.6 0.01 0.001 0.04 0.003 0.001 752 854 * REM in steel grade M is Y: 0.001, Ce: 0.001.
 Tables 6 to 10 lead to the following conclusions. The steel sheets of Tests No. 1, 2, 4, 5, 10, 11, 13 to 16, 18 to 23, 25 to 28, and 30 to 35 which were produced using the steel grades A to N and Q to U having the component composition in accordance with the present invention and under the annealing conditions stipulated by the present invention had excellent ductility and bendability in a tensile strength range of 980 MPa and higher.
 The amount of C in steel grade O and the amount of Mn in steel grade P in Table 6 exceeded the respective upper limits of the present invention and fracture occurred during cold rolling. As a result, sample steels could not be produced.
 Test No. 3 is an example in which the time of holding at the reheating holding temperature after cooling to 500.degree. C. or less was short. The amount of MA structure increased and the bendability was degraded.
 Test No. 6 is an example in which no holding at the predetermined temperature was performed after coiling. The standard deviation of the regions with the concentration of Mn of 1.2 times was low and the bendability was degraded.
 Test No. 7 is an example in which the coiling temperature and holding temperature were low. The area ratio of the region with the concentration of Mn of 1.2 times and the standard deviation of the regions with the concentration of Mn of 1.2 times were low and the bendability was degraded.
 Test No. 8 is an example in which the time of holding at the predetermined temperature after coiling was short. The standard deviation of the regions with the concentration of Mn of 1.2 times was low and the bendability was degraded.
 Test No, 9 is an example in which the reheating holding temperature after cooling to 500.degree. C. or less was low. The amount of MA structure increased and the bendability was degraded.
 Test No. 12 is an example in which the soaking temperature was high. Ferrite was generated insufficiently and the concentration of Mn in the ferrite phase was too high. As a result, the ductility was degraded.
 Test No. 17 is an example in which the cooling stop temperature was high. The amount of ferrite and MA structure increased, the strength was low, and the bendability degraded.
 Test No, 24 is an example in which the cooling rate to 500.degree. C. was low. The concentration of Mn in the ferrite phase became too high. As a result, the bendability was degraded.
 Test No. 29 is an example in which no reheating was performed after cooling to 500.degree. C. or less. The amount of the MA structure increased and the bendability was degraded.

References: Application No. 2014
 Application No. 2014
 Application No. 2014
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 Application No. 2014
 Application No. 2014