STEEL SHEET, MEMBER, AND METHODS FOR MANUFACTURING SAME

Abstract

Provided are a steel sheet; a related member; and methods for manufacturing the same. The steel sheet has a chemical composition including, in mass %, C: 0.06 to 0.25%, Si: 0.4 to 2.5%, Mn: 1.5 to 3.5%, P: 0.02% or less, S: 0.01% or less, sol. Al: less than 1.0%, and N: less than 0.015%, the balance being Fe and incidental impurities, the steel sheet being such that the steel sheet includes a steel microstructure including, in area fraction, polygonal ferrite: 10% or less (including 0%), tempered martensite: 40% or more, fresh martensite: 20% or less (including 0%), bainitic ferrite having 20 or less internal carbides per 10 m.sup.2: 3 to 40%, and, in volume fraction, retained austenite: 5 to 20%, and the steel sheet has S.sub.C0.5/S.sub.C0.3100 of 20% or more.

Claims

1-11. (canceled)

12. A steel sheet having a chemical composition comprising, in mass %: C: 0.06 to 0.25%, Si: 0.4 to 2.5%, Mn: 1.5 to 3.5%, P: 0.02% or less, S: 0.01% or less, sol. Al: less than 1.0%, N: less than 0.015%, and optionally, one or two or more selected from the following groups A to C: group A: one or two selected from: Ti: 0.1% or less, and B: 0.01% or less, group B: one, or two or more selected from: Cu: 1% or less, Ni: 1% or less, Cr: 1.0% or less, Mo: 0.5% or less, V: 0.5% or less, Nb: 0.1% or less, Zr: 0.2% or less, and W: 0.2% or less, group C: one, or two or more selected from: Ca: 0.0040% or less, Ce: 0.0040% or less, La: 0.0040% or less, Mg: 0.0030% or less, Sb: 0.1% or less, and Sn: 0.1% or less, the balance being Fe and incidental impurities, the steel sheet being such that: the steel sheet comprises a steel microstructure including, in area fraction, polygonal ferrite: 10% or less (including 0%), tempered martensite: 40% or more, fresh martensite: 20% or less (including 0%), bainitic ferrite having 20 or less internal carbides per 10 m.sup.2: 3 to 40%, and, in volume fraction, retained austenite: 5 to 20%, and the steel sheet has a proportion S.sub.C0.5/S.sub.C0.3100 of 20% or more wherein S.sub.C0.5 is the area of a region having a C concentration of 0.50% or more and S.sub.C0.3 is the area of a region having a C concentration of 0.30% or more.

13. The steel sheet according to claim 12, wherein the steel microstructure has a number density of retained austenite present adjacent to the bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 of 50 or more per 10000 m.sup.2.

14. The steel sheet according to claim 12, which has a galvanized layer on a surface.

15. The steel sheet according to claim 13, which has a galvanized layer on a surface.

16. A member obtained using the steel sheet described in claim 12.

17. A member obtained using the steel sheet described in claim 13.

18. A member obtained using the steel sheet described in claim 14.

19. A member obtained using the steel sheet described in claim 15.

20. A method for manufacturing a steel sheet, comprising hot rolling and cold rolling a steel slab having the chemical composition described in claim 12, and annealing the cold rolled steel sheet, the annealing comprising: a step of holding the steel sheet at an annealing temperature of 810 to 900 C.; a step of cooling the steel sheet in a range of temperatures from 810 C. to 500 C. at an average cooling rate CR1 of 5 to 100 C./s; a step of causing the steel sheet to reside in a range of temperatures from 500 C. to a residence finish temperature T1 that is equal to or higher than a martensite start temperature Ms ( C.) and is equal to or higher than 320 C. for 10 seconds or more and 60 seconds or less while cooling the steel sheet at an average cooling rate CR2 of 10 C./s or less; a step of cooling the steel sheet in a range of temperatures from the residence finish temperature T1 to a finish cooling temperature T2 of 200 C. or above and 300 C. or below at an average cooling rate CR3 of 3 to 100 C./s; a step of heating the steel sheet in a range of temperatures from the finish cooling temperature T2 to 380 C. at an average heating rate of 2 C./s or more; a step of causing the steel sheet to reside in a range of temperatures of 340 C. or above and 590 C. or below for 20 seconds or more and 3000 seconds or less while cooling the steel sheet at an average cooling rate CR4 of 0.01 to 5 C./s; and a step of cooling the steel sheet to a temperature of 50 C. or below at an average cooling rate CR5 of 0.1 C./s or more.

21. The method for manufacturing a steel sheet according to claim 20, wherein the step of causing the steel sheet to reside at an average cooling rate CR4 of 0.01 to 5 C./s includes performing a hot-dip galvanizing treatment or a hot-dip galvannealing treatment.

22. The method for manufacturing a steel sheet according to claim 20, further comprising a step of performing an electrogalvanizing treatment after the step of cooling the steel sheet at an average cooling rate CR5 of 0.1 C./s or more.

23. A method for manufacturing a member, comprising a step of subjecting the steel sheet described in claim 12 to at least one working of forming and joining to produce a member.

24. A method for manufacturing a member, comprising a step of subjecting the steel sheet described in claim 13 to at least one working of forming and joining to produce a member.

25. A method for manufacturing a member, comprising a step of subjecting the steel sheet described in claim 14 to at least one working of forming and joining to produce a member.

26. A method for manufacturing a member, comprising a step of subjecting the steel sheet described in claim 15 to at least one working of forming and joining to produce a member.

Description

BRIEF DESCRIPTION OF THE DRAWINGS

[0096] FIG. 1 is a set of views illustrating a manner for evaluating the laser weldability of a steel sheet according to aspects of the present invention.

[0097] FIG. 2 is an example SEM image of a steel microstructure of a steel sheet.

[0098] FIG. 3 is a set of views illustrating a manner for measuring a steel microstructure of a steel sheet according to aspects of the present invention.

[0099] FIG. 4 is a diagram illustrating a method for manufacturing a steel sheet according to aspects of the present invention.

DETAILED DESCRIPTION OF EMBODIMENTS OF THE INVENTION

[0100] Embodiments of the present invention will be described in detail below. The present invention is not limited to the embodiments described below.

[0101] A steel sheet according to aspects of the present invention has a chemical composition including, in mass %, C: 0.06 to 0.25%, Si: 0.4 to 2.5%, Mn: 1.5 to 3.5%, P: 0.02% or less, S: 0.01% or less, sol. Al: less than 1.0%, and N: less than 0.015%, the balance being Fe and incidental impurities. The steel sheet includes a steel microstructure including, in area fraction, polygonal ferrite: 10% or less (including 0%), tempered martensite: 40% or more, fresh martensite: 20% or less (including 0%), bainitic ferrite having 20 or less internal carbides per 10 m.sup.2: 3 to 40%, and, in volume fraction, retained austenite: 5 to 20%. The steel sheet has a proportion S.sub.C0.5/S.sub.C0.3100 of 20% or more wherein S.sub.C0.5 is the area of a region having a C concentration of 0.50% or more and S.sub.C0.3 is the area of a region having a C concentration of 0.30% or more.

[0102] The steel sheet according to aspects of the present invention will be described below in the order of its chemical composition and its steel microstructure.

[0103] The steel sheet according to aspects of the present invention includes the components described below. In the following description, the unit % for the contents of components means mass %.

C: 0.06 to 0.25%

[0104] Carbon is added to ensure predetermined strength by ensuring an area fraction of tempered martensite, to enhance ductility by ensuring a volume fraction of retained , and to stabilize retained by being concentrated in retained and thereby to enhance ductility. Furthermore, the addition of carbon increases the strength of a fused portion of a welded joint and a portion quenched from region, and thereby can eliminate or reduce deformation occurring at HAZ and enhance the HAZ softening resistance. When the C content is less than 0.06%, these effects cannot be ensured sufficiently. Thus, the lower limit is limited to 0.06%. The C content is preferably 0.09% or more, and more preferably 0.11% or more. When the C content exceeds 0.25%, upper bainite transformation during intermediate holding in the course of cooling is retarded, and it becomes difficult to form a predetermined amount of retained that is adjacent to upper bainite. As a result, ductility is lowered. Furthermore, the amount of massive martensite or massive retained is increased to deteriorate stretch flange formability. Furthermore, laser welding characteristics of the steel sheet, such as HAZ softening resistance, spot weldability, bendability, and flangeability, are significantly deteriorated. Thus, the upper limit of the C content is limited to 0.25%. From the points of view of ductility and HAZ softening resistance, the C content is preferably 0.22% or less. In order to further improve ductility and HAZ softening resistance, the C content is more preferably 0.20% or less.

Si: 0.4 to 2.5%

[0105] Silicon is added to offer high strength by strengthening ferrite, to enhance the stability of retained by suppressing the formation of carbides in martensite and bainite and thereby to enhance ductility, and to enhance the HAZ softening resistance of a weld by increasing the amount of solid solution strengthening that is insusceptible to heat. From the points of view of the above, the Si content is limited to 0.4% or more. In order to enhance ductility, the Si content is preferably 0.6% or more. More preferably, the Si content is 0.8% or more. When the Si content exceeds 2.5%, the rolling load at the time of hot rolling is extremely increased to make sheet production difficult. Furthermore, chemical convertibility and weld toughness are deteriorated. For these reasons, the Si content is limited to 2.5% or less. To ensure chemical convertibility and the toughness of the base material and a weld, the Si content is preferably less than 2.0%. To ensure weld toughness, the Si content is preferably 1.8% or less, and more preferably 1.5% or less.

Mn: 1.5 to 3.5%

[0106] Manganese is an important element from the points of view of ensuring strength by ensuring a predetermined area fraction of tempered martensite and/or bainite; improving ductility by lowering the Ms temperature of retained and thereby stabilizing retained ; enhancing ductility by suppressing the formation of carbides in bainite similarly to silicon; and enhancing ductility by increasing the volume fraction of retained . In order to obtain these effects, the Mn content is limited to 1.5% or more. In order to enhance ductility by stabilizing retained , the Mn content is preferably 2.5% or more. The Mn content is preferably 2.6% or more, and more preferably 2.7% or more. When the Mn content exceeds 3.5%, bainite transformation is significantly retarded to lower ductility and HAZ softening resistance. When the Mn content exceeds 3.5%, moreover, it becomes difficult to suppress the formation of massive coarse and massive coarse martensite, resulting in a decrease in stretch flange formability. Thus, the Mn content is limited to 3.5% or less. In order to ensure high ductility by promoting bainite transformation, the Mn content is preferably 3.2% or less. More preferably, the Mn content is 3.1% or less.

P: 0.02% or Less

[0107] Phosphorus is an element that strengthens steel, but much phosphorus deteriorates spot weldability. Thus, the P content is limited to 0.02% or less. In order to improve spot weldability, the P content is preferably 0.01% or less. The P content may be nil. From the point of view of manufacturing cost, the P content is preferably 0.001% or more.

S: 0.01% or Less

[0108] Sulfur is an element that is effective in improving scale exfoliation in hot rolling and effective in suppressing nitridation during annealing, but sulfur lowers spot weldability, bendability, and flangeability. From the points of view of the above, the S content is limited to 0.01% or less. In accordance with aspects of the present invention, the contents of C, Si, and Mn are high and spot weldability tends to be lowered. In order to improve spot weldability, the S content is preferably 0.0020% or less, and more preferably less than 0.0010%. The S content may be nil. From the point of view of manufacturing cost, the S content is preferably 0.0001% or more. More preferably, the S content is 0.0005% or more.

Sol. Al: Less than 1.0%

[0109] Aluminum is added for the purposes of deoxidization and stabilizing retained as a substitute for silicon. The lower limit of the sol. Al content is not particularly limited. For stable deoxidization, the sol. Al content is preferably 0.005% or more. The sol. Al content is more preferably 0.01% or more. On the other hand, 1.0% or more sol. Al significantly lowers the strength of the base material and also affects adversely chemical convertibility. Thus, the sol. Al content is limited to less than 1.0%. In order to obtain high strength, the sol. Al content is more preferably less than 0.50%, and still more preferably 0.20% or less.

N: Less than 0.015%

[0110] Nitrogen is an element that forms nitrides, such as BN, AlN, and TiN, in steel. This element lowers the hot ductility of steel and lowers the surface quality. Furthermore, in B-containing steel, nitrogen has a harmful effect in eliminating the effect of boron through the formation of BN. The surface quality is significantly deteriorated when the N content is 0.015% or more. Thus, the N content is limited to less than 0.015%. The N content is preferably 0.010% or less. The N content may be nil. From the point of view of manufacturing cost, the N content is preferably 0.0001% or more. More preferably, the N content is 0.001% or more.

[0111] The balance after the above components is Fe and incidental impurities. The steel sheet according to aspects of the present invention preferably has a chemical composition that contains the basic components described above, with the balance consisting of iron (Fe) and incidental impurities.

[0112] In addition to the above components, the chemical composition of the steel sheet according to aspects of the present invention may appropriately include one, or two or more optional elements selected from the following (A), (B), and (C): [0113] (A) in mass %, one or two selected from Ti: 0.1% or less and B: 0.01% or less, [0114] (B) in mass %, one, or two or more selected from Cu: 1% or less, Ni: 1% or less, Cr: 1.0% or less, Mo: 0.5% or less, V: 0.5% or less, Nb: 0.1% or less, Zr: 0.2% or less, and W: 0.2% or less, [0115] (C) in mass %, one, or two or more selected from Ca: 0.0040% or less, Ce: 0.0040% or less, La: 0.0040% or less, Mg: 0.0030% or less, Sb: 0.1% or less, and Sn: 0.1% or less.

Ti: 0.1% or Less

[0116] Titanium fixes nitrogen in steel as TiN to produce an effect of enhancing hot ductility and an effect of allowing boron to produce its effect of enhancing hardenability. Furthermore, titanium has an effect of reducing the size of the microstructure through TiC precipitation. In order to obtain these effects, the Ti content is preferably 0.002% or more. In order to fix nitrogen sufficiently, the Ti content is more preferably 0.008% or more. The Ti content is still more preferably 0.010% or more. On the other hand, more than 0.1% titanium may cause an increase in rolling load and a decrease in ductility by an increased amount of precipitation strengthening. Thus, when titanium is added, the Ti content is limited to 0.1% or less. Preferably, the Ti content is 0.05% or less. In order to ensure high ductility, the Ti content is more preferably 0.03% or less.

B: 0.01% or Less

[0117] Boron is an element that enhances the hardenability of steel and facilitates the formation of a predetermined area fraction of tempered martensite and/or bainite. Furthermore, boron enhances the hardenability in the vicinity of a weld and allows a hard microstructure to be formed in the vicinity of the weld, thereby enhancing HAZ softening resistance. Furthermore, residual solute boron enhances delayed fracture resistance. In order to obtain these effects of boron, the B content is preferably 0.0002% or more. The B content is more preferably 0.0005% or more. Still more preferably, the B content is 0.0010% or more. When, on the other hand, the B content exceeds 0.01%, the effects are saturated, and further hot ductility is significantly lowered to invite surface defects. Thus, when boron is added, the B content is limited to 0.01% or less. Preferably, the B content is 0.0050% or less. More preferably, the B content is 0.0030% or less.

Cu: 1% or Less

[0118] Copper enhances the corrosion resistance in automobile use environments. Furthermore, corrosion products of copper cover the surface of the steel sheet and can suppress penetration of hydrogen into the steel sheet. Copper is an element that is mixed when scraps are used as raw materials. By accepting copper contamination, recycled materials can be used as raw materials and thereby manufacturing costs can be reduced. From these points of view, the Cu content is preferably 0.005% or more, and, further from the point of view of enhancing delayed fracture resistance, the Cu content is more preferably 0.05% or more. Still more preferably, the Cu content is 0.10% or more. On the other hand, too much copper invites surface defects. Thus, when copper is added, the Cu content is limited to 1% or less. The Cu content is preferably 0.4% or less, and more preferably 0.2% or less.

Ni: 1% or Less

[0119] Similar to copper, nickel can enhance corrosion resistance. Furthermore, nickel can also eliminate or reduce the occurrence of surface defects that tend to occur when the steel contains copper. To benefit from these effects, it is preferable to add 0.01% or more nickel. The Ni content is more preferably 0.04% or more, and still more preferably 0.06% or more. On the other hand, adding too much nickel can instead cause surface defects because scales are formed nonuniformly in a heating furnace, and also increases the cost. Thus, when nickel is added, the Ni content is limited to 1% or less. The Ni content is preferably 0.4% or less, and more preferably 0.2% or less.

Cr: 1.0% or Less

[0120] Chromium may be added to produce an effect of enhancing the hardenability of steel and an effect of suppressing the formation of carbides in martensite and upper/lower bainite. Furthermore, chromium enhances the hardenability in the vicinity of a weld and allows a hard phase to be formed in the vicinity of the weld, thereby enhancing HAZ softening resistance. In order to obtain these effects, the Cr content is preferably 0.01% or more. The Cr content is more preferably 0.03% or more, and still more preferably 0.06% or more. On the other hand, too much chromium deteriorates pitting corrosion resistance. Thus, when chromium is added, the Cr content is limited to 1.0% or less. The Cr content is preferably 0.8% or less, and more preferably 0.4% or less.

Mo: 0.5% or Less

[0121] Molybdenum may be added to produce an effect of enhancing the hardenability of steel and an effect of suppressing the formation of carbides in martensite and upper/lower bainite. Furthermore, molybdenum enhances the hardenability in the vicinity of a weld and allows a hard phase to be formed in the vicinity of the weld, thereby enhancing HAZ softening resistance. In order to obtain these effects, the Mo content is preferably 0.01% or more. The Mo content is more preferably 0.03% or more, and still more preferably 0.06% or more. On the other hand, molybdenum significantly deteriorates the chemical convertibility of the cold rolled steel sheet. Thus, when molybdenum is added, the Mo content is limited to 0.5% or less. From the point of view of enhancing chemical convertibility, the Mo content is preferably 0.15% or less.

V: 0.5% or Less

[0122] Vanadium may be added to produce an effect of enhancing the hardenability of steel, an effect of suppressing the formation of carbides in martensite and upper/lower bainite, an effect of reducing the size of the microstructure, and an effect of improving delayed fracture resistance through the precipitation of carbide. Furthermore, vanadium enhances the hardenability in the vicinity of a weld and allows a hard phase to be formed in the vicinity of the weld, thereby enhancing HAZ softening resistance. In order to obtain these effects, the V content is preferably 0.003% or more. The V content is more preferably 0.005% or more, and still more preferably 0.010% or more. On the other hand, much vanadium significantly deteriorates castability. Thus, when vanadium is added, the V content is limited to 0.5% or less. The V content is preferably 0.3% or less, and more preferably 0.1% or less. The V content is still more preferably 0.05% or less, and further preferably 0.03% or less.

Nb: 0.1% or Less

[0123] Niobium may be added to produce an effect of reducing the size of the steel microstructure and thereby increasing the strength, and, through grain size reduction, an effect of promoting bainite transformation, an effect of improving bendability, and an effect of enhancing delayed fracture resistance. Furthermore, niobium enhances the hardenability in the vicinity of a weld and allows a hard phase to be formed in the vicinity of the weld, thereby enhancing HAZ softening resistance. In order to obtain these effects, the Nb content is preferably 0.002% or more. The Nb content is more preferably 0.004% or more, and still more preferably 0.010% or more. On the other hand, adding much niobium results in excessive precipitation strengthening and low ductility. Furthermore, the rolling load is increased and castability is deteriorated. Thus, when niobium is added, the Nb content is limited to 0.1% or less. The Nb content is preferably 0.05% or less, and more preferably 0.03% or less.

Zr: 0.2% or Less

[0124] Zirconium may be added to produce an effect of enhancing the hardenability of steel, an effect of suppressing the formation of carbides in bainite, an effect of reducing the size of the microstructure, and an effect of improving delayed fracture resistance through the precipitation of carbide. In order to obtain these effects, the Zr content is preferably 0.005% or more. The Zr content is more preferably 0.008% or more, and still more preferably 0.010% or more. When, on the other hand, the steel contains much zirconium, increased amounts of coarse precipitates, such as ZrN and ZrS, remain undissolved at the time of slab heating before hot rolling to cause deterioration in delayed fracture resistance. Thus, when zirconium is added, the Zr content is limited to 0.2% or less. The Zr content is preferably 0.15% or less, and more preferably 0.08% or less. The Zr content is still more preferably 0.03% or less, and further preferably 0.02% or less.

W: 0.2% or Less

[0125] Tungsten may be added to produce an effect of enhancing the hardenability of steel, an effect of suppressing the formation of carbides in bainite, an effect of reducing the size of the microstructure, and an effect of improving delayed fracture resistance through the precipitation of carbide. In order to obtain these effects, the W content is preferably 0.005% or more. The W content is more preferably 0.008% or more, and still more preferably 0.010% or more.

[0126] When, on the other hand, the steel contains much tungsten, increased amounts of coarse precipitates, such as WN and WS, remain undissolved at the time of slab heating before hot rolling to cause deterioration in delayed fracture resistance. Thus, when tungsten is added, the W content is limited to 0.2% or less. The W content is preferably 0.15% or less, and more preferably 0.08% or less. The W content is still more preferably 0.03% or less, and further preferably 0.02% or less.

Ca: 0.0040% or Less

[0127] Calcium fixes sulfur as CaS and contributes to improvements in bendability and delayed fracture resistance. Thus, the Ca content is preferably 0.0002% or more. The Ca content is more preferably 0.0005% or more, and still more preferably 0.0010% or more. On the other hand, much calcium deteriorates surface quality and bendability. Thus, when calcium is added, the Ca content is limited to 0.0040% or less. The Ca content is preferably 0.0035% or less, and more preferably 0.0020% or less.

Ce: 0.0040% or Less

[0128] Similar to calcium, cerium fixes sulfur and contributes to improvements in bendability and delayed fracture resistance. Thus, the Ce content is preferably 0.0002% or more. The Ce content is more preferably 0.0004% or more, and still more preferably 0.0006% or more. On the other hand, much cerium deteriorates surface quality and bendability. Thus, when cerium is added, the Ce content is limited to 0.0040% or less. The Ce content is preferably 0.0035% or less, and more preferably 0.0020% or less.

La: 0.0040% or Less

[0129] Similar to calcium, lanthanum fixes sulfur and contributes to improvements in bendability and delayed fracture resistance. Thus, the La content is preferably 0.0002% or more. The La content is more preferably 0.0004% or more, and still more preferably 0.0006% or more. On the other hand, much lanthanum deteriorates surface quality and bendability. Thus, when lanthanum is added, the La content is limited to 0.0040% or less. The La content is preferably 0.0035% or less, and more preferably 0.0020% or less.

Mg: 0.0030% or Less

[0130] Magnesium fixes oxygen as MgO and contributes to improvement in delayed fracture resistance. Thus, the Mg content is preferably 0.0002% or more. The Mg content is more preferably 0.0004% or more, and still more preferably 0.0006% or more. On the other hand, much magnesium deteriorates surface quality and bendability. Thus, when magnesium is added, the Mg content is limited to 0.0030% or less. The Mg content is preferably 0.0025% or less, and more preferably 0.0010% or less.

Sb: 0.1% or Less

[0131] Antimony suppresses oxidation and nitridation of a superficial portion of the steel sheet and thereby eliminates or reduces the loss of the C and B contents in the superficial portion. Furthermore, the elimination or reduction of the loss of the C and B contents leads to suppressed formation of ferrite in the superficial portion of the steel sheet, thus increasing strength and improving delayed fracture resistance. From these points of view, the Sb content is preferably 0.002% or more. The Sb content is more preferably 0.004% or more, and still more preferably 0.006% or more. When, on the other hand, the Sb content exceeds 0.1%, castability is deteriorated and segregation occurs at prior grain boundaries to deteriorate the delayed fracture resistance of sheared end faces. Thus, when antimony is added, the Sb content is limited to 0.1% or less. The Sb content is preferably 0.04% or less, and more preferably 0.03% or less.

Sn: 0.1% or Less

[0132] Tin suppresses oxidation and nitridation of a superficial portion of the steel sheet and thereby eliminates or reduces the loss of the C and B contents in the superficial portion. Furthermore, the elimination or reduction of the loss of the C and B contents leads to suppressed formation of ferrite in the superficial portion of the steel sheet, thus increasing strength and improving delayed fracture resistance. From these points of view, the Sn content is preferably 0.002% or more. The Sn content is preferably 0.004% or more, and more preferably 0.006% or more. When, on the other hand, the Sn content exceeds 0.1%, castability is deteriorated. Furthermore, tin is segregated at prior grain boundaries to deteriorate the delayed fracture resistance of sheared end faces. Thus, when tin is added, the Sn content is limited to 0.1% or less. The Sn content is preferably 0.04% or less, and more preferably 0.03% or less.

[0133] When the content of any of the above optional components is below the preferred lower limit, the optional element present below the lower limit does not impair the advantageous effects according to aspects of the present invention. Thus, such an optional element below the preferred lower limit content is regarded as an incidental impurity.

[0134] Next, the steel microstructure of the steel sheet according to aspects of the present invention will be described.

[0135] Polygonal ferrite: 10% or less (including 0%) Polygonal ferrite, which is formed during annealing or a cooling process, contributes to enhancement in ductility but decreases stretch flange formability by giving rise to a difference in hardness from surrounding hard phases, such as martensite. Polygonal ferrite does not impair the advantageous effects according to aspects of the present invention as long as the area fraction thereof is 10% or less and therefore, polygonal ferrite may be contained up to 10% or less in area fraction. Thus, in accordance with aspects of the present invention, the area fraction of polygonal ferrite is limited to 10% or less. The polygonal ferrite is preferably 5% or less, and more preferably 2% or less. The polygonal ferrite may be 0%.

Tempered Martensite: 40% or More

[0136] In order to obtain predetermined strength and stretch flange formability, the area fraction of tempered martensite is limited to 40% or more. The tempered martensite is preferably 50% or more. When, on the other hand, the tempered martensite exceeds 80%, the strength is excessively increased and the ductility is lowered. Thus, the tempered martensite is preferably 80% or less. The tempered martensite is more preferably 75% or less.

Fresh Martensite: 20% or Less (Including 0%)

[0137] The final tempering step (the residence step at an average cooling rate CR4 described later) to produce a large amount of bainite transformation conventionally results in a large amount of massive martensite or massive retained that remains. The conventional approach to preventing this is to reduce the amount of manganese and thereby to promote bainite transformation. However, decreasing the Mn content lowers ductility because of the loss of the effects of stabilizing retained and increasing the volume fraction of retained . In contrast, aspects of the present invention perform an appropriate cooling treatment on the steel sheet containing a large amount of manganese, and thereby can make use of bainite transformation and reduce the occurrence of massive microstructures at the same time.

[0138] Excellent stretch flange formability and HAZ softening resistance can be ensured by reducing the area fraction of massive fresh martensite microstructures to 20% or less. Thus, the area fraction of fresh martensite in accordance with aspects of the present invention is limited to 20% or less. In order to ensure excellent stretch flange formability and HAZ softening resistance, the fresh martensite is preferably 10% or less. The fresh martensite is more preferably 5% or less. The fresh martensite may be 0%.

Bainitic Ferrite Having 20 or Less Internal Carbides Per 10 m.SUP.2.: 3 to 40%

[0139] When bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 are present with an area fraction of 3% or more, carbon is efficiently concentrated into surrounding retained . Furthermore, such bainitic ferrite is insusceptible to heat at the time of laser welding and contributes to enhancement in HAZ softening resistance. Thus, in accordance with aspects of the present invention, the area fraction of bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 is limited to 3% or more. The bainitic ferrite is preferably 5% or more, and more preferably 7% or more. In order to reduce the decrease in strength, the area fraction of bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 is limited to 40% or less.

[0140] The area fraction is preferably 30% or less, and more preferably 25% or less.

[0141] In accordance with aspects of the present invention, the steel may contain bainitic ferrite having more than 20 internal carbides per 10 m.sup.2.

Microstructure Including One, or Two or More of Tempered Martensite, Fresh Martensite, Upper Bainite, Lower Bainite, and Retained Austenite: 90% or More (Including 100%)

[0142] In order to ensure predetermined strength, ductility, and stretch flange formability, it is preferable that the remaining microstructure after the polygonal ferrite have a total area fraction of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite of 90% or more. The remaining microstructure may be a microstructure including one, or two or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite, or may be a microstructure consisting of one, or two or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite.

[0143] The upper bainite and the lower bainite include bainitic ferrite having 20 or less internal carbides per 10 m.sup.2. The upper bainite and the lower bainite may include bainitic ferrite having more than 20 internal carbides per 10 m.sup.2.

Retained Austenite: 5 to 20%

[0144] In order to ensure high ductility, the volume fraction of retained austenite (retained ) is limited to 5% or more of the entire steel microstructure. The retained austenite is preferably 7% or more, and more preferably 9% or more. This amount of retained includes retained formed adjacent to bainite. An excessively large amount of retained invites decreases in strength, stretch flange formability, and delayed fracture resistance. Thus, the volume fraction of retained is limited to 20% or less. The retained austenite is preferably 15% or less. Incidentally, the volume fraction may be regarded as the area fraction.

Proportion S.sub.C0.5/S.sub.C0.3100 of 20% or More where S.sub.C0.5 is the Area of a Region Having a C Concentration of 0.50% (Mass %) or More and S.sub.C0.3 is the Area of a Region Having a C Concentration of 0.30% (Mass %) or More

[0145] In order to ensure high ductility, the proportion S.sub.C0.5/S.sub.C0.3100 is limited to 20% or more. Here, S.sub.C0.5 is the area of a region having a C concentration of 0.50% or more and S.sub.C0.3 is the area of a region having a C concentration of 0.30% or more. The proportion is preferably 25% or more, and more preferably 30% or more.

(Preferred Requirement) Number Density of Retained Austenite Present Adjacent to Bainitic Ferrite Having 20 or Less Internal Carbides Per 10 m.sup.2: 50 or More Per 10000 m.sup.2

[0146] In the formation of bainitic ferrite having 20 or less internal carbides per 10 m.sup.2, carbon is efficiently partitioned to adjacent non-transformed austenite, and consequently retained in the final microstructure attains a high carbon concentration and can contribute to enhancement in ductility. From the point of view of ensuring higher ductility, it is preferable in accordance with aspects of the present invention that the number density of retained austenite present adjacent to bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 be 50 or more per 10000 m.sup.2. More preferably, the number density is 70 or more, and still more preferably 100 or more per 10000 m.sup.2. To avoid a decrease in strength due to excessive formation of retained , the number density of retained austenite present adjacent to bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 is preferably 400 or less, and more preferably 300 or less per 10000 m.sup.2.

[0147] Next, a method for measuring the steel microstructure of the steel sheet according to aspects of the present invention will be described.

[0148] To measure the area fractions of polygonal ferrite, bainitic ferrite, tempered martensite, and fresh martensite, the steel sheet is cut to expose a through-thickness cross section that is perpendicular to the steel sheet surface and is parallel to the rolling direction. The cross section is mirror-polished and is etched with 3 vol % Nital. Portions at thickness are observed with SEM in 10 fields of view at a magnification of 5000 times. FIG. 2 illustrates an example SEM image of the steel microstructure of the steel sheet. As illustrated in FIG. 2, the polygonal ferrite discussed here is relatively equiaxed ferrite containing almost no internal carbides. This region looks blackest in SEM. Bainitic ferrite is a ferrite microstructure that contains internal carbide or retained which looks white in SEM.

[0149] When determination is difficult whether the ferrite is bainitic ferrite or polygonal ferrite, the area fractions are calculated while assuming that the region is polygonal ferrite when the aspect ratio is 2.0 and the region is bainitic ferrite when the aspect ratio is >2.0.

[0150] FIG. 3 is a set of views illustrating a manner for measuring the steel microstructure of the steel sheet according to aspects of the present invention. As illustrated in FIG. 3(A), the aspect ratio is obtained from a/b in which a is the major axis length where the particle length is longest, and b is the minor axis length that cuts across the particle over the largest length perpendicular to the major axis length. When particles are in contact with one another, as illustrated in FIG. 3(B), the particles are divided at a position that divides the particles approximately evenly and the sizes of the respective particles are measured.

[0151] The number of internal carbides in bainitic ferrite per 10 m.sup.2 can be determined by measuring the area of bainitic ferrite and counting the number of internal carbides in a 5000 SEM image, dividing the carbide count by the area of bainitic ferrite, and converting the quotient into the value per 10 m.sup.2.

[0152] Tempered martensite is a region that contains a lath-like submicrostructure and carbide precipitates according to SEM. Fresh martensite is a massive region that looks white and does not contain any visible submicrostructures according to SEM.

[0153] The microstructure including one, or two or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite corresponds to the remaining microstructure after the polygonal ferrite, and the total area fraction of this microstructure is the area fraction of the regions other than the polygonal ferrite. Here, the area fraction of carbides is very small and is thus included in the above area fraction of the remaining microstructure.

[0154] The volume fraction of retained austenite (retained ) was determined by chemically polishing the steel sheet surface to a location at thickness and analyzing the sheet surface by X-ray diffractometry. Co-K radiation source is used as the incident X-ray, and the volume fraction of retained austenite is calculated from the intensity ratio of (200), (211), and (220) planes of ferrite and (200), (220), and (311) planes of austenite. Because retained is randomly distributed, the volume fraction of retained obtained by X-ray diffractometry is equal to the area fraction of retained in the steel microstructure.

[0155] The number density of retained austenite present adjacent to bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 is determined as follows. The sample used for bainitic ferrite observation is mirror-polished. An electron backscattering diffraction pattern (EBSD) of the same field of view as obtained in SEM is subjected to mapping measurement using EBSD analysis program OIM Data Collection ver. 7. The obtained data is analyzed using TSL OIM Analysis ver. 7 (manufactured by EDAX/TSL) to give phase map data. The number density of fcc structures present adjacent to bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 is measured. Here, the term adjacent means that the fcc structures are in contact with bcc structures corresponding to bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 in the phase map. The fcc structure may be included in the bcc structure.

[0156] The area S.sub.C0.5 of a region having a C concentration of 0.50% (mass %) or more and the area S.sub.C0.3 of a region having a C concentration of 0.30% (mass %) or more are measured by mapping analysis of the C concentration distribution with respect to positions at thickness of a through-thickness cross section perpendicular to the steel sheet surface and parallel to the rolling direction, using field emission electron probe microanalyzer (FE-EPMA) JXA-8500F manufactured by JEOL Ltd., at an acceleration voltage of 6 kV and an illumination current of 710.sup.8 A with the minimum beam diameter. In order to eliminate the influence of contamination, the background is subtracted so that the average value of carbon obtained by the analysis will be equal to the amount of carbon in the base material. Specifically, when the average of the measured amounts of carbon is greater than the amount of carbon in the base material, the excess is understood as contamination, and the excess is subtracted from each of the values analyzed at the respective positions. The values thus obtained are taken as the true amounts of carbon at the respective positions.

[0157] The steel sheet according to aspects of the present invention preferably has a tensile strength of 980 MPa or more. More preferably, the tensile strength is 1180 MPa or more. The upper limit of the tensile strength is preferably 1450 MPa or less from the point of view of compatibility with other characteristics, and is more preferably 1400 MPa or less.

[0158] In the steel sheet according to aspects of the present invention, the total elongation T-El is 16.0% or more when TS is less than 1180 MPa, 14.0% or more when TS is 1180 MPa or more and less than 1320 MPa, and 13.0% or more when TS is 1320 MPa or more. With this configuration, forming stability is markedly enhanced. It is preferable to ensure that the hole expansion ratio is 30% or more. From the point of view of compatibility with other characteristics, the upper limit of A is preferably 90% or less, and more preferably 80% or less at any level of strength.

[0159] The steel sheet according to aspects of the present invention is preferably such that a specimen obtained by laser welding is fractured in the base material fracture mode in a fracture mode determination test and satisfies HAZ strength base material TS+50 MPa in a notched tensile test.

[0160] The steel sheet according to aspects of the present invention described above may be a steel sheet having a galvanized layer on a surface. The galvanized layer may be a hot-dip galvanized layer or an electrogalvanized layer.

[0161] Next, a method for manufacturing the steel sheet according to aspects of the present invention will be described.

[0162] In a method for manufacturing the steel sheet according to aspects of the present invention, a steel slab having the chemical composition described hereinabove is hot rolled and cold rolled. The cold rolled steel sheet obtained is annealed. The annealing includes the following steps in the order named: a step of holding the steel sheet at an annealing temperature of 810 to 900 C.; a step of cooling the steel sheet in a range of temperatures from 810 C. to 500 C. at an average cooling rate (CR1) of 5 to 100 C./s; a step of causing the steel sheet to reside in a range of temperatures from 500 C. to a residence finish temperature (T1) that is equal to or higher than a martensite start temperature Ms ( C.) and is equal to or higher than 320 C. for 10 seconds or more and 60 seconds or less while cooling the steel sheet at an average cooling rate (CR2) of 10 C./s or less; a step of cooling the steel sheet in a range of temperatures from the residence finish temperature (T1) to a finish cooling temperature (T2) of 200 C. or above and 300 C. or below at an average cooling rate (CR3) of 3 to 100 C./s; a step of heating the steel sheet in a range of temperatures from the finish cooling temperature (T2) to 380 C. at an average heating rate of 2 C./s or more; a step of causing the steel sheet to reside in a range of temperatures of 340 C. or above and 590 C. or below for 20 seconds or more and 3000 seconds or less while cooling the steel sheet at an average cooling rate (CR4) of 0.01 to 5 C./s; and a step of cooling the steel sheet to a temperature of 50 C. or below at an average cooling rate (CR5) of 0.1 C./s or more.

[0163] The temperatures specified in the steps in accordance with aspects of the present invention indicate the surface temperatures of the slab (steel slab) or the steel sheet.

[0164] FIG. 4 is a diagram illustrating the method for manufacturing the steel sheet according to aspects of the present invention, in particular, indicating changes in surface temperature of the slab (steel slab) or the steel sheet with time. The details of the steps, including the changes in temperature with time, will be described below.

Hot Rolling

[0165] For example, the steel slab may be hot rolled in such a manner that the slab is heated and then rolled, that the slab from continuous casting is subjected to hot direct rolling without heating, or that the slab from continuous casting is quickly heat treated and then rolled. The hot rolling may be performed in accordance with a conventional procedure. For example, the slab heating temperature may be 1100 to 1300 C.; the soaking time may be 20 to 300 minutes; the finish rolling temperature may be Ar.sub.3 transformation temperature to Ar.sub.3 transformation temperature+200 C.; and the coiling temperature may be 400 to 720 C. In order to eliminate or reduce thickness variations and to ensure high strength stably, the coiling temperature is preferably 430 to 530 C.

Cold Rolling

[0166] In cold rolling, the rolling reduction ratio (the cumulative rolling reduction ratio) may be 30 to 85%. In order to ensure high strength stably and to reduce anisotropy, the rolling reduction ratio is preferably 35 to 85%. When a high rolling load is incurred, a softening annealing treatment may be performed on CAL (a continuous annealing line) or in BAF (a box annealing furnace) at 450 to 730 C.

Annealing

[0167] After the steel slab having the aforementioned chemical composition is hot rolled and cold rolled, the steel sheet is annealed under the conditions specified below. The annealing facility is not particularly limited, but a continuous annealing line (CAL) or a continuous galvanizing line (CGL) is preferable from the points of view of productivity and ensuring the desired heating rate and cooling rate.

Holding at an Annealing Temperature of 810 to 900 C.

[0168] In order to ensure the predetermined area fraction of tempered martensite and/or bainite and to ensure the predetermined volume fraction of retained , the annealing temperature is limited to 810 to 900 C. In order to control polygonal ferrite to 5% or less, the annealing temperature is preferably adjusted so that the annealing will take place in the single-phase region. The annealing temperature is preferably 815 C. or above. When, on the other hand, the annealing temperature exceeds 900 C., the grain size is excessively increased to extend the distance of diffusion of carbon atoms required to obtain retained having the desired carbon concentration, resulting in a decrease in ductility. Thus, the annealing temperature is limited to 900 C. or below. Preferably, the annealing temperature is 880 C. or below.

Cooling in a Range of Temperatures from 810 C. to 500 C. at an Average Cooling Rate (CR1) of 5 to 100 C./s

[0169] After the steel sheet is held at 810 to 900 C., the steel sheet is cooled in a range of temperatures from 810 C. to 500 C. at an average cooling rate (CR1) of 5 to 100 C./s. When the average cooling rate (CR1) is lower than 5 C./s, a large amount of ferrite is formed to cause a decrease in strength and a decrease in stretch flange formability. The average cooling rate (CR1) is preferably 8 C./s or more. When, on the other hand, the average cooling rate (CR1) is too high, the sheet shape is deteriorated. Thus, the average cooling rate is limited to 100 C./s or less. The average cooling rate (CR1) is preferably 50 C./s or less, and more preferably less than 30 C./s.

[0170] Here, the average cooling rate (CR1) is (810 C. (cooling start temperature)500 C. (finish cooling temperature))/(cooling time (seconds) from cooling start temperature 810 C. to finish cooling temperature 500 C.).

Residence in a Range of Temperatures from 500 C. to a Residence Finish Temperature (T1) that is Equal to or Higher than a Martensite Start Temperature Ms ( C.) and is Equal to or Higher than 320 C. for 10 Seconds or More and 60 Seconds or Less while Performing Cooling at an Average Cooling Rate (CR2) of 10 C./s or Less

[0171] The steel sheet is caused to reside (is gradually cooled) in a range of temperatures from 500 C. to a residence finish temperature (T1) that is equal to or higher than a martensite start temperature Ms ( C.) and is equal to or higher than 320 C. for 10 seconds or more and 60 seconds or less while being cooled at an average cooling rate (CR2) of 10 C./s or less. In this manner, bainite having a low carbide density can be formed, and retained having a high C concentration can be formed adjacent thereto. When the temperature range is below the Ms or below 320 C., martensite transformation precedes and lower bainite is formed, thus causing a decrease in strength. When, on the other hand, the temperature range exceeds 500 C., the driving force for bainite transformation decreases and the amount of bainite transformation is reduced. Thus, the temperature range is limited to be 500 C. or below and to be equal to or higher than the Ms and equal to or higher than 320 C. This temperature range is preferably 380 C. or above, and more preferably 420 C. or above. The temperature range is preferably 480 C. or below, and more preferably 460 C. or below. When the average cooling rate (CR2) exceeds 10 C./s, the amount of bainite transformation is reduced. Thus, the average cooling rate (CR2) is limited to 10 C./s or less. When the residence time is less than 10 seconds, the desired amount of bainite cannot be obtained. When the residence time exceeds 60 seconds, the enrichment of carbon from bainite to massive non-transformed proceeds to result in an increase in the residual amount of the massive microstructure. Thus, the residence time is limited to 10 seconds or more and 60 seconds or less. In order to ensure bainitic ferrite and retained austenite and to enhance ductility and HAZ softening resistance, the residence time is preferably 20 seconds or more. In order to reduce the formation of massive microstructures and thereby to enhance stretch flange formability, the residence time is preferably 50 seconds or less.

[0172] The martensite start temperature Ms can be determined using a Formaster tester by holding a cylindrical test piece (3 mm in diameter10 mm in height) at a predetermined annealing temperature and quenching the test piece with helium gas while measuring the volume change.

[0173] Here, the average cooling rate (CR2) is (500 C. (residence start temperature)residence finish temperature (T1))/(residence time (seconds) from 500 C. to residence finish temperature (T1)).

Cooling in a Range of Temperatures from the Residence Finish Temperature (T1) to a Finish Cooling Temperature (T2) of 200 C. or Above and 300 C. or Below at an Average Cooling Rate (CR3) of 3 to 100 C./s

[0174] After the above residence, the steel sheet needs to be cooled rapidly to avoid excessive progress of carbon enrichment into . When the average cooling rate (CR3) in the range of temperatures from the residence finish temperature T1 of 320 C. or above to the finish cooling temperature T2 of 200 C. or above and 300 C. or below is less than 3 C./s, carbon is concentrated into massive non-transformed and an increased amount of fresh martensite is formed during the final cooling to cause a decrease in stretch flange formability. From the point of view of enhancing stretch flange formability, the average cooling rate (CR3) in the range of temperatures from the residence finish temperature T1 to the finish cooling temperature T2 of 200 C. or above and 300 C. or below is limited to 3 C./s or more. The average cooling rate (CR3) is more preferably 5 C./s or more, and still more preferably 8 C./s or more. When the average cooling rate in the above temperature range is excessively high, the sheet shape is deteriorated. Thus, the average cooling rate (CR3) in the above temperature range is limited to 100 C./s or less. The average cooling rate (CR3) is preferably 50 C./s or less. In order to ensure a predetermined amount of retained , the finish cooling temperature T2 is limited to 200 C. or above. The finish cooling temperature T2 is preferably 220 C. or above, and more preferably 240 C. or above. When the finish cooling temperature T2 is above 300 C., a large amount of massive non-transformed remains and an increased amount of fresh martensite is formed during the final cooling to cause a decrease in stretch flange formability. Thus, the finish cooling temperature T2 is limited to 300 C. or below. The finish cooling temperature T2 is preferably 280 C. or below.

[0175] Here, the average cooling rate (CR3) is (residence finish temperature (T1))(finish cooling temperature (T2))/(cooling time (seconds) from residence finish temperature (T1) to finish cooling temperature (T2)).

Heating in a Range of Temperatures from the Finish Cooling Temperature (T2) to 380 C. at an Average Heating Rate of 2 C./s or More

[0176] Furthermore, the steel sheet is heated in a range of temperatures from the finish cooling temperature (T2) to 380 C. in a short time. In this manner, the precipitation of carbides is suppressed and high ductility can be ensured. When the steel sheet is reheated to 380 C. or above, upper bainite is formed from martensite or bainite as nucleus that has been formed during cooling. When the average heating rate during heating to 380 C. is low, the above effects cannot be obtained. As a result, the amount of retained is reduced and ductility is lowered. Thus, the average heating rate in the range of temperatures from the finish cooling temperature (T2) to 380 C. is limited to 2 C./s or more. From the points of view of suppressing carbide precipitation and forming upper bainite at the time of reheating, the average heating rate is preferably 5 C./s or more, and more preferably 10 C./s or more. The upper limit of the average heating rate is not particularly limited but is preferably 50 C./s or less, and more preferably 30 C./s or less.

[0177] Here, the average heating rate is 380 C. (heating stop temperature)(finish cooling temperature (T2))/(heating time (seconds) from finish cooling temperature T2 to 380 C. (heating stop temperature)).

Residence in a Range of Temperatures of 340 C. or Above and 590 C. or Below for 20 Seconds or More and 3000 Seconds or Less while Performing Cooling at an Average Cooling Rate (CR4) of 0.01 to 5 C./s

[0178] In order to partition carbon to retained and stabilize the retained and in order to divide massive regions distributed as non-transformed by bainite transformation and thereby to enhance stretch flange formability, the steel sheet is caused to reside (is gradually cooled) in a range of temperatures of 340 C. or above and 590 C. or below for 20 seconds or more and 3000 seconds or less. Furthermore, in order to enhance stretch flange formability by eliminating or reducing the formation of massive microstructures due to excessive partitioning of carbon to retained and also by causing self-tempering of fresh martensite, the steel sheet is cooled slowly in this temperature range at an average cooling rate (CR4) of 0.01 to 5 C./s. When the average cooling rate (CR4) is less than 0.01 C./s, carbon is excessively partitioned to retained and massive microstructures are formed to cause a decrease in stretch flange formability. Thus, the average cooling rate (CR4) is limited to 0.01 C./s or more. When, on the other hand, the average cooling rate (CR4) exceeds 5 C./s, the partitioning of carbon to retained is suppressed and a sufficient amount of carbon-enriched regions cannot be obtained. Furthermore, fresh martensite is formed to cause a decrease in A. Thus, the average cooling rate (CR4) is limited to 5 C./s or less.

[0179] Here, the average cooling rate (CR4) is (cooling start temperature (T3))(finish cooling temperature (T4))/(cooling time (seconds) from cooling start temperature (T3) to finish cooling temperature (T4)).

[0180] Here, the cooling start temperature (T3) and the finish cooling temperature (T4) are not particularly limited as long as they are in the range of 340 C. or above and 590 C. or below. The cooling start temperature (T3) is preferably in the range of 360 to 580 C. The finish cooling temperature (T4) is preferably in the range of 350 to 450 C.

[0181] The holding (residence) in the temperature range of 340 to 590 C. may also include a hot-dip galvanizing treatment. That is, the steel sheet may be subjected to a hot-dip galvanizing treatment or a hot-dip galvannealing treatment in the step where the steel sheet is caused to reside while being cooled at an average cooling rate (CR4) of 0.01 to 5 C./s. When a hot-dip galvanizing treatment is performed, the steel sheet is hot-dip galvanized by being immersed into a galvanizing bath at 440 C. or above and 500 C. or below, and the coating weight is preferably adjusted by, for example, gas wiping. The galvanizing bath used in the hot-dip galvanization preferably has an Al content of 0.10% or more and 0.22% or less. Furthermore, a hot-dip galvannealing treatment may be performed by an alloying treatment of the zinc coating after the hot-dip galvanizing treatment. The alloying treatment of the zinc coating is preferably performed in a temperature range of 470 C. or above and 590 C. or below. Although this step is a cooling step (residence and slow cooling), the hot-dip galvanizing treatment and the alloying treatment of the zinc coating may be performed during the step as long as the temperature range, the residence time, and the average cooling rate CR4 described above are satisfied. The hot-dip galvanizing treatment and the alloying treatment of the zinc coating may involve a temperature rise.

Cooling to a Temperature of 50 C. or Below at an Average Cooling Rate (CR5) of 0.1 C./s or More

[0182] In order to prevent softening due to excessive tempering and to avoid a decrease in ductility due to carbide precipitation, the steel sheet is then cooled to a temperature of 50 C. or below at an average cooling rate (CR5) of 0.1 C./s or more. In order to stabilize press formability, for example, to control surface roughness and to flatten the sheet shape, and also in order to increase the YS, the steel sheet may be subjected to skin pass rolling. The skin pass rolling reduction ratio is preferably 0.1 to 0.5%. The sheet shape may be flattened with a leveler. The average cooling rate (CR5) to a temperature of 50 C. or below is preferably 5 C./s or more, and more preferably 100 C./s or less.

[0183] Here, the average cooling rate (CR5) is (340 C. (cooling start temperature)finish cooling temperature of 50 C. or below)/(cooling time (seconds) from cooling start temperature to finish cooling temperature).

[0184] In order to improve stretch flange formability, the above heat treatment or the skin pass rolling may be followed by a low-temperature heat treatment at 100 to 300 C. for 30 seconds to 10 days. This treatment tempers martensite that has been formed during the final cooling or the skin pass rolling and detaches from the steel sheet hydrogen that has penetrated into the steel sheet during annealing. By the low-temperature heat treatment, hydrogen can be reduced to less than 0.1 ppm. Furthermore, electroplating may be performed. That is, the steel sheet may be electrogalvanized after the step where the steel sheet is cooled at an average cooling rate (CR5) of 0.1 C./s or more. The electrogalvanization is preferably followed by the above low-temperature heat treatment in order to reduce the amount of hydrogen in the steel.

[0185] The steel sheet according to aspects of the present invention preferably has a thickness of 0.5 mm or more. The thickness is preferably 2.0 mm or less.

[0186] Next, a member and a method for manufacture thereof according to aspects of the present invention will be described.

[0187] The member according to aspects of the present invention is obtained by subjecting the steel sheet according to aspects of the present invention to at least one working of forming and joining. The method for manufacturing a member according to aspects of the present invention includes a step of subjecting the steel sheet according to aspects of the present invention to at least one working of forming and joining to produce a member.

[0188] The steel sheet according to aspects of the present invention has a tensile strength of 980 MPa or more and is excellent in ductility, stretch flange formability, and laser weldability. Thus, the member that is obtained using the steel sheet according to aspects of the present invention also has high strength and has excellent ductility, excellent stretch flange formability, and excellent laser weldability compared to the conventional high-strength members. Furthermore, weight can be reduced by using the member according to aspects of the present invention. Thus, for example, the member according to aspects of the present invention may be suitably used in an automobile body frame part. The member according to aspects of the present invention also includes a welded joint.

[0189] The forming may be performed using any common working process, such as press working, without limitation. Furthermore, the joining may be performed using common welding, such as spot welding or arc welding, or, for example, riveting or crimping without limitation.

EXAMPLES

[0190] EXAMPLES of the present invention will be described below.

[0191] Steel sheets according to aspects of the present invention and steel sheets of COMPARATIVE EXAMPLES were manufactured by treating 1.4 mm thick cold rolled steel sheets having a chemical composition described in Table 1 under annealing conditions described in Table 2.

[0192] The cold rolled steel sheets had been obtained by subjecting steel slabs having the chemical composition described in Table 1 to hot rolling (slab heating temperature: 1200 C., soaking time: 60 minutes, finish rolling temperature: 900 C., coiling temperature: 500 C.) and cold rolling (rolling reduction ratio (cumulative rolling reduction ratio): 50%).

[0193] In Table 2, the martensite start temperature Ms was obtained using a Formaster tester by holding a cylindrical test piece (3 mm in diameter10 mm in height) at a predetermined annealing temperature and quenching the test piece with helium gas while measuring the volume change.

[0194] Some of the steel sheets (cold rolled steel sheets: CR) were obtained as hot-dip galvanized steel sheets (GI) by being hot-dip galvanized in a step where the steel sheet was caused to reside in a range of temperatures of 340 C. or above and 590 C. or below for 20 seconds or more and 3000 seconds or less while being cooled at an average cooling rate of 0.01 to 5 C./s. Here, the steel sheets were hot-dip galvanized by being immersed into a galvanizing bath at a temperature of 440 C. or above and 500 C. or below, and the coating weight was adjusted by, for example, gas wiping. The galvanizing bath used in the hot-dip galvanization had an Al content of 0.10% or more and 0.22% or less. Furthermore, some of the hot-dip galvanized steel sheets were alloyed and obtained as hot-dip galvannealed steel sheets (GA) by being subjected to an alloying treatment after the hot-dip galvanizing treatment. Here, the alloying treatment was performed in a range of temperatures of 460 C. or above and 590 C. or below. Furthermore, some of the steel sheets (cold rolled steel sheets: CR) were obtained as electrogalvanized steel sheets (EG) by electroplating.

[0195] The steel microstructure was measured in the following manner. The measurement results are described in Table 3. To measure the area fractions of polygonal ferrite, bainitic ferrite, tempered martensite, and fresh martensite, the steel sheet was cut to expose a through-thickness cross section that was parallel to the rolling direction. The cross section was mirror-polished and was etched with 3 vol % Nital. Portions at thickness were observed with SEM in 10 fields of view at a magnification of 5000 times. As illustrated in FIG. 2, the polygonal ferrite discussed here is relatively equiaxed ferrite containing almost no internal carbides. This region looks blackest in SEM. Bainitic ferrite is a ferrite microstructure that contains carbide or retained which looks white in SEM.

[0196] When determination was difficult whether the ferrite was bainitic ferrite or polygonal ferrite, the area fractions were calculated while assuming that the region was polygonal ferrite when the aspect ratio was 2.0 and the region was bainitic ferrite when the aspect ratio was >2.0.

[0197] As illustrated in FIG. 3(A), the aspect ratio was obtained from a/b in which a was the major axis length where the particle length was longest, and b was the minor axis length that cut across the particle over the largest length perpendicular to the major axis length. When particles were in contact with one another, as illustrated in FIG. 3(B), the particles were divided at a position that divided the particles approximately evenly and the sizes of the respective particles were measured.

[0198] The number of internal carbides in bainitic ferrite per 10 m.sup.2 was determined by measuring the area of bainitic ferrite and counting the number of internal carbides in a 5000 SEM image, dividing the carbide count by the area of bainitic ferrite, and converting the quotient into the value per 10 m.sup.2.

[0199] Tempered martensite is a region that contains a lath-like submicrostructure and carbide precipitates according to SEM. Fresh martensite is a massive region that looks white and does not contain any visible submicrostructures according to SEM.

[0200] The microstructure including one, or two or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite corresponds to the remaining microstructure after the polygonal ferrite, and the total area fraction of this microstructure is the area fraction of the regions other than the polygonal ferrite. Here, the area fraction of carbides was very small and was thus included in the above area fraction of the remaining microstructure.

[0201] The volume fraction of retained austenite (retained ) was determined by chemically polishing the steel sheet surface to a location at thickness and analyzing the sheet surface by X-ray diffractometry. Co-K radiation source was used as the incident X-ray, and the volume fraction of retained austenite was calculated from the intensity ratio of (200), (211), and (220) planes of ferrite and (200), (220), and (311) planes of austenite. Because retained is randomly distributed, the volume fraction of retained obtained by X-ray diffractometry is equal to the area fraction of retained in the steel microstructure.

[0202] The number density of retained austenite present adjacent to bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 was determined as follows. The sample used for bainitic ferrite observation was mirror-polished. An electron backscattering diffraction pattern (EBSD) of the same field of view as obtained in SEM was subjected to mapping measurement using EBSD analysis program OIM Data Collection ver. 7. The obtained data was analyzed using TSL OIM Analysis ver. 7 (manufactured by EDAX/TSL) to give phase map data. The number density of fcc structures present adjacent to bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 was measured. Here, the term adjacent means that the fcc structures were in contact with bcc structures corresponding to bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 in the phase map. The fcc structure may be included in the bcc structure.

[0203] The area S.sub.C0.5 of a region having a C concentration of 0.50% (mass %) or more and the area S.sub.C0.3 of a region having a C concentration of 0.30% (mass %) or more were measured by mapping analysis of the C concentration distribution with respect to positions at thickness of a through-thickness cross section perpendicular to the steel sheet surface and parallel to the rolling direction, using field emission electron probe microanalyzer (FE-EPMA) JXA-8500F manufactured by JEOL Ltd., at an acceleration voltage of 6 kV and an illumination current of 710.sup.8 A with the minimum beam diameter.

[0204] In order to eliminate the influence of contamination, the background was subtracted so that the average value of carbon obtained by the analysis would be equal to the amount of carbon in the base material. Specifically, when the average of the measured amounts of carbon was greater than the amount of carbon in the base material, the excess was understood as contamination, and the excess was subtracted from each of the values analyzed at the respective positions. The values thus obtained were taken as the true amounts of carbon at the respective positions.

[0205] JIS No. 5 test pieces for tensile test and test pieces for hole expansion test were sampled from the steel sheets obtained. A tensile test was performed (in accordance with JIS Z2241). The tensile strength TS and the total elongation T-El are described in Table 3. The steel sheets were evaluated as being excellent in strength when the tensile strength was 980 MPa or more. Furthermore, the ductility was evaluated as excellent when the total elongation T-El was 16.0% or more for the steel sheets having a TS of less than 1180 MPa, when the total elongation T-El was 14.0% or more for the steel sheets having a TS of 1180 MPa or more and less than 1320 MPa, and when the total elongation T-El was 13.0% or more for the steel sheets having a TS of 1320 MPa or more.

[0206] Furthermore, the test pieces for hole expansion test sampled from the steel sheets after the heat treatment were subjected to a hole expansion test conforming to the provisions of The Japan Iron and Steel Federation Standard JFST 1001 to evaluate stretch flange formability.

[0207] Specifically, a 100 mm100 mm square sample was punched with a punching tool having a punch diameter of 10 mm and a die diameter of 10.3 mm (13% clearance), and a conical punch having an apex angle of 60 degrees was inserted into the hole in such a manner that the burr produced at the time of punching would be directed to the outside. The hole was expanded until the sheet was cracked through the thickness. The hole expansion ratio (%)={(dd.sub.0)/d.sub.0}100 was calculated. Here, d.sub.0: initial hole diameter (mm), and d: hole diameter (mm) at the occurrence of cracking. The results are described in Table 3. The steel was evaluated as having excellent flangeability when A was 30% or more.

[0208] Furthermore, two steel sheet workpieces each 120 mm in the direction perpendicular to the rolling direction and 200 mm in the rolling direction were sampled from the steel sheet (the end faces were ground). The two workpieces were butted against each other so that their rolling directions were aligned. The butt position was laser welded. The gap between the butted faces was 0 mm. The laser welding was performed using Nd-YAG laser. The spot diameter at the focus position was 0.6 mm. The focus position was 4 mm above the steel sheet. The shield gas was Ar. The laser output was 4.2 kW. The welding speed was 3.7 m/min. A test piece for tensile test having a shape described in FIG. 1(a) was sampled from the welded member in such a manner that the weld line would be perpendicular to the tensile axis and would be located at the longitudinal center of the test piece. The test piece was subjected to a tensile test to determine the fracture mode (fracture mode determination test). When the fracture position was 2.0 mm or more away from the weld line (when even a part of the fracture position was 2.0 mm or more away from the weld line), the fracture mode was determined to be the base material fracture. When the fracture position was less than 2.0 mm away from the weld line and the test piece had been fractured by the extension of a crack along the weld line (the extension of a crack in the HAZ or the fused portion), the fracture mode was determined to be the weld fracture. Furthermore, a test piece was sampled from the welded member in such a manner that the weld line would be perpendicular to the tensile axis and would be located at the longitudinal center of the test piece, and the weld was notched as illustrated in FIG. 1(b). The notched test piece was subjected to a tensile test (a notched tensile test). In the manner described above, the test pieces were deformed exclusively at the HAZ of the weld and limited regions around the HAZ and were forcibly fractured at the HAZ portion to evaluate the strength of the HAZ portion itself. Thus, more quantitative grasp of the quality of laser weldability was realized. In the evaluation, the laser weldability (the HAZ softening resistance) was evaluated as being excellent when the test piece was fractured in the base material fracture mode in the fracture mode determination test and satisfied HAZ strength base material TS+50 MPa in the notched tensile test.

[0209] INVENTIVE EXAMPLES described in Tables 2 and 3 attained excellent strength, ductility, flangeability, and laser weldability (HAZ softening resistance). In contrast, COMPARATIVE EXAMPLES were unsatisfactory in one or more of these properties.

TABLE-US-00001 TABLE 1 Chemical composition (mass %) Steel C Si Mn P S sol. Al N others Remarks A 0.125 0.63 3.04 0.004 0.0005 0.043 0.0030 Compliant steel B 0.264 1.17 2.32 0.009 0.0010 0.017 0.0071 Comparative steel C 0.113 1.24 2.62 0.001 0.0010 0.029 0.0042 Ti: 0.027, B: 0.0031 Compliant steel D 0.134 0.34 2.43 0.004 0.0005 10.032 0.0024 Comparative steel E 0.198 1.79 2.51 0.003 0.0006 0.039 0.0056 Ti: 0.018, B: 0.0054, Nb: 0.030 Compliant steel F 0.201 1.68 1.42 0.002 0.0010 0.031 0.0030 Comparative steel G 0.241 2.32 1.58 0.002 0.0018 0.039 0.0012 Ti: 0.042, B: 0.0028, Cu: 0.18, Compliant Ni: 0.06, Cr: 0.06, Mo: 0.06 steel H 0.177 1.61 2.93 0.003 0.0012 0.009 0.0056 V: 0.013, Zr: 0.010, W: 0.011 Compliant steel I 0.048 1.37 2.64 0.005 0.0005 0.006 0.0053 Comparative steel J 0.186 0.62 3.30 0.010 0.0006 0.005 0.0078 Ti: 0.021, B: 0.0050, Compliant Ca: 0.0012, Ce: 0.0007, steel La: 0.0020 K 0.208 1.95 2.71 0.001 0.0004 0.006 0.0067 B: 0.0016, Mg: 0.0010, Compliant Sb: 0.01, Sn: 0.01 steel L 0.164 2.59 3.41 0.007 0.0006 0.040 0.0056 Comparative steel M 0.181 2.43 3.61 0.006 0.0005 0.021 0.0063 Comparative steel *Underlines indicate being outside of the range of the present invention. *The balance after the above components is Fe and incidental impurities.

TABLE-US-00002 TABLE 2 Annealing conditions Finish Residence cooling Residence finish temp. Heating Annealing CR1*1 CR2*2 time*3 temp. T1 CR3*4 T2 rate*5 No. Steel temp. ( C.) ( C./s) Ms( C.) ( C./s) (sec) ( C.) ( C./s) ( C.) ( C./s) 1 A 820 30 330 1 40 460 3 220 2 2 A 875 20 355 1 60 440 5 280 5 3 A 800 30 310 1 60 440 5 250 5 4 A 910 30 355 1 60 440 5 250 5 5 B 830 30 330 1 60 440 5 260 5 6 C 845 6 370 3 10 470 8 260 10 7 C 815 10 345 3 20 440 10 205 10 8 C 830 3 330 3 20 440 10 260 10 9 C 830 10 340 15 10 350 10 260 10 10 D 830 10 375 3 20 440 10 220 10 11 E 830 20 375 3 30 410 10 250 10 12 F 860 20 375 3 30 410 10 295 20 13 E 860 20 375 3 8 476 10 220 20 14 E 800 20 330 2 80 340 10 280 20 15 F 820 20 345 3 50 350 10 280 20 16 G 810 20 330 3 50 350 10 230 20 17 G 840 20 345 3 50 350 10 260 20 18 G 840 20 345 3 50 350 1 280 20 19 G 840 20 345 3 50 350 10 190 20 20 H 890 50 335 5 10 450 10 280 20 21 H 890 50 335 5 30 350 20 280 20 22 H 840 50 330 5 30 350 20 310 20 23 H 840 50 330 5 30 350 20 290 1 24 I 850 50 380 3 30 410 20 250 20 25 J 850 60 325 5 30 350 30 260 20 26 J 850 60 325 5 30 350 30 260 20 27 J 815 60 315 5 30 350 30 280 20 28 J 850 60 325 5 30 350 30 220 20 29 K 810 80 310 10 10 400 50 210 20 30 K 810 95 310 10 18 320 80 240 20 31 K 860 95 325 10 10 400 80 230 20 32 K 830 95 325 10 10 400 80 280 20 33 L 830 95 300 10 10 400 80 290 20 34 M 840 95 300 10 10 400 80 290 20 Annealing conditions Cooling Finish Residence start cooling time temp. temp. CR4*6 *7 T3 T4 CR5*8 No. Steel ( C./s) (sec) ( C.) ( C.) ( C./s) Plating*9 Remarks 1 A 0.01 2800 380 352 5 CR INV. EX. 2 A 0.03 1500 420 375 5 CR INV. EX. 3 A 0.03 1500 420 375 5 CR COMP EX. 4 A 0.03 1500 420 375 5 CR COMP. EX. 5 B 0.03 1500 420 375 5 CR COMP. EX. 6 C 0.03 500 400 385 5 CR INV. EX. 7 C 0.03 500 400 385 5 CR INV. EX. 8 C 0.03 500 400 385 5 CR COMP. EX. 9 C 0.03 500 400 385 5 CR COMP. EX. 10 D 0.03 500 400 385 5 CR COMP EX. 11 E 0.05 700 450 415 5 CR INV. EX. 12 F 0.05 700 450 415 5 CR INV. EX. 13 E 0.05 700 450 415 5 CR COMP EX. 14 E 0.05 700 450 415 5 CR COMP EX. 15 F 0.05 700 450 415 5 CR COMP. EX. 16 G 0.07 900 550 487 5 GA INV. EX. 17 G 0.07 900 550 487 5 GA INV. EX. 18 G 0.07 900 550 487 5 GA COMP EX. 19 G 0.07 900 550 487 5 GA COMP. EX. 20 H 0.20 600 500 380 80 GI INV. EX. 21 H 0.20 600 500 380 10 GI INV. EX. 22 H 0.20 600 500 380 10 GI COMP. EX. 23 H 0.20 600 500 380 10 GI COMP. EX. 24 I 0.20 600 500 380 10 GI COMP. EX. 25 C 0.80 300 585 345 10 EG INV. EX. 26 J 0.80 300 585 345 10 EG INV. EX. 27 J 0.00 100 350 350 10 EG COMP. EX. 28 J 10 30 480 180 10 EG COMP. EX. 29 K 1.50 100 500 350 1 CR INV. EX. 30 K 3.00 50 500 350 1 CR INV. EX. 31 K 4.00 10 430 390 1 EG COMP. EX. 32 K 0.01 3200 390 358 1 EG COMP. EX 33 L 0.07 900 440 377 1 EG COMP. EX. 34 M 0.07 1000 440 370 1 EG COMP. EX. *Underlines indicate being outside of the range of the present invention. *1Average cooling rate CR1 in range of temperatures from 810 to 500 C. *2Average cooling rate CR2 in range of temperatures from 500 C. to residence finish temperature T1 *3Residence time in range of temperatures from 500 C. to residence finish temperature T1 *4Average cooling rate CR3 in range of temperatures from residence finish temperature T1 to finish cooling temperature T2 *5Average heating rate in range of temperatures from finish cooling temperature T2 to 380 C. *6Average cooling rate CR4 in range of temperatures of 340 C. or above and 590 C. or below *7 Residence time in range of temperatures of 340 C. or above and 590 C. or below *8Average cooling rate CR5 to temperature of 50 C. or below *9CR: no plating, GA: hot-dip galvannealed steel sheet, GI: hot-dip galvanized steel sheet (no alloying treatment of zinc coating), EG: electrogalvanized steel sheet

TABLE-US-00003 TABLE 3 Microstructure Tempered Fresh .sub.BF*11 Remaining Polygonal martensite martensite BF*10 number micro- Retained y ferrite area area area area density structure*12 volume fraction fraction fraction fraction (precipitates/ area fraction fraction No. Steel (%) (%) (%) (%) 10000 m.sup.2) (%) (%) 1 A 9 41 8 34 230 91 8 2 A 0 55 6 30 202 100 9 3 A 13 43 22 16 148 87 6 4 A 0 59 3 31 205 100 7 5 B 1 63 18 8 213 99 10 6 C 4 66 7 15 174 96 8 7 C 8 67 8 7 142 92 10 8 C 12 54 15 11 153 88 8 9 C 5 75 9 2 83 95 9 10 D 4 73 6 13 52 96 4 11 E 1 74 5 6 130 99 12 12 E 0 72 13 9 148 100 6 13 E 1 87 3 2 66 99 7 14 E 5 24 24 41 242 95 6 15 F 15 56 17 8 163 85 4 16 G 3 74 4 5 147 97 14 17 G 2 67 6 8 137 98 12 18 G 2 26 21 43 235 98 8 19 G 1 78 3 15 188 99 3 20 H 0 72 11 6 72 100 11 21 H 0 69 9 12 122 100 10 22 H 1 63 23 11 135 99 2 23 H 2 74 17 4 140 98 3 24 I 45 34 5 15 32 55 1 25 J 1 68 7 13 150 99 11 26 J 0 61 7 15 161 100 12 27 J 8 51 22 9 243 92 10 28 J 0 80 4 10 55 100 6 29 K 8 66 4 4 48 92 18 30 K 7 58 8 7 63 93 15 31 K 0 81 3 12 45 100 4 32 K 4 41 11 41 174 96 3 33 L 2 62 23 2 163 98 11 34 M 1 62 22 3 163 99 12 Characteristics Microstructure HAZ S.sub.C0.5/ TS T-El Fracture strength No. Steel S.sub.C0.3(%) (MPa) (%) (%) position (MPa) Remarks 1 A 41 991 18.8 53 Base 1052 INV. EX. material 2 A 40 1003 18.1 55 Base 1067 INV. EX. material 3 A 38 975 18.4 28 HAZ 1010 COMP. EX. 4 A 18 1108 15.8 64 Base 1166 COMP. EX. material 5 B 48 1421 5.3 26 Base 1482 COMP. EX. material 6 C 37 1067 17.5 65 Base 1132 INV. EX. material 7 C 39 1082 16.7 62 Base 1141 INV. EX. material 8 C 32 968 19.3 45 Base 1031 COMP. EX. material 9 C 17 1056 15.6 53 Base 1102 COMP. EX. material 10 D 21 1069 15.1 56 Base 1126 COMP. EX. material 11 E 36 1201 17.2 51 Base 1256 INV. EX. material 12 E 22 1268 15.3 47 Base 1330 INV. EX. material 13 E 18 1287 13.7 62 Base 1335 COMP. EX. material 14 E 40 1183 16.5 27 HAZ 1221 COMP. EX. 15 F 24 1154 13.6 32 Base 1210 COMP. EX. material 16 G 39 1333 14.7 58 Base 1394 INV. EX. material 17 G 36 1350 14.2 57 Base 1407 INV. EX. material 18 G 38 1325 15.1 24 HAZ 1361 COMP. EX. 19 G 35 1407 12.8 42 Base 1463 COMP. EX. material 20 H 31 1258 16.2 53 Base 1316 INV. EX. material 21 H 30 1241 16.8 59 Base 1297 INV. EX. material 22 H 26 1285 13.5 26 HAZ 1322 COMP. EX. 23 H 29 1263 13.8 33 Base 1325 COMP. EX. material 24 I 15 554 18.0 58 Base 611 COMP. EX. material 25 J 35 1237 16.5 55 Base 1293 INV. EX. material 26 J 36 1230 17.4 51 Base 1286 INV. EX. material 27 J 38 1214 15.9 28 HAZ 1251 COMP. EX. 28 J 17 1296 13.4 59 Base 1356 COMP. EX. material 29 K 23 1222 15.1 48 Base 1279 INV. EX. material 30 K 28 1215 16.2 43 Base 1275 INV. EX. material 31 K 16 1313 13.2 63 Base 1373 COMP. EX. material 32 K 38 1224 13.8 42 Base 1288 COMP. EX. material 33 L 36 1288 13.4 28 HAZ 1330 COMP. EX. 34 M 37 1293 13.1 28 HAZ 1340 COMP. EX. *Underlines indicate being outside of the range of the present invention. *10Bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 *11Retained austenite present adjacent to bainitic ferrite having 20 or less internal carbides per 10 m.sup.2 *12Microstructure including one, or two or more of tempered martensite, fresh martensite, upper bainite, lower bainite, and retained austenite

[0210] The steel sheets according to aspects of the present invention have superior ductility, excellent stretch flange formability, and excellent laser weldability and can be suitably applied to press forming and be suitably used in the press forming process in the manufacturing of, for example, automobiles and home appliances.

[0211] The steel sheets of INVENTIVE EXAMPLES have high strength, excellent ductility, excellent stretch flange formability, and excellent laser weldability. This has shown that members obtained by forming of the steel sheets of INVENTIVE EXAMPLES, members obtained by joining of the steel sheets of INVENTIVE EXAMPLES, and members obtained by forming and joining of the steel sheets of INVENTIVE EXAMPLES will have high strength, excellent ductility, excellent stretch flange formability, and excellent laser weldability similar to the steel sheets of INVENTIVE EXAMPLES.