EP2695961B1 - Tôle d'acier à haute résistance présentant une excellente aptitude au façonnage, et son procédé de fabrication - Google Patents

Tôle d'acier à haute résistance présentant une excellente aptitude au façonnage, et son procédé de fabrication Download PDF

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EP2695961B1
EP2695961B1 EP12765664.3A EP12765664A EP2695961B1 EP 2695961 B1 EP2695961 B1 EP 2695961B1 EP 12765664 A EP12765664 A EP 12765664A EP 2695961 B1 EP2695961 B1 EP 2695961B1
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Prior art keywords
temperature
steel sheet
content
temperature range
bainite
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EP12765664.3A
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German (de)
English (en)
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EP2695961A4 (fr
EP2695961A1 (fr
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Yuichi Futamura
Michiharu Nakaya
Takayuki Kimura
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Kobe Steel Ltd
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Kobe Steel Ltd
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Priority claimed from JP2011197671A external-priority patent/JP5685167B2/ja
Priority claimed from JP2011197670A external-priority patent/JP5685166B2/ja
Application filed by Kobe Steel Ltd filed Critical Kobe Steel Ltd
Priority to EP15172192.5A priority Critical patent/EP2942416B1/fr
Publication of EP2695961A1 publication Critical patent/EP2695961A1/fr
Publication of EP2695961A4 publication Critical patent/EP2695961A4/fr
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    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/34—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the shape of the material to be treated
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
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    • C21D6/00—Heat treatment of ferrous alloys
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    • C21D6/00—Heat treatment of ferrous alloys
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
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    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
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    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
    • Y10T—TECHNICAL SUBJECTS COVERED BY FORMER US CLASSIFICATION
    • Y10T428/00—Stock material or miscellaneous articles
    • Y10T428/12—All metal or with adjacent metals
    • Y10T428/12493—Composite; i.e., plural, adjacent, spatially distinct metal components [e.g., layers, joint, etc.]
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    • Y10T428/12785—Group IIB metal-base component
    • Y10T428/12792—Zn-base component
    • Y10T428/12799—Next to Fe-base component [e.g., galvanized]

Definitions

  • the present invention relates to high-strength steel sheets having excellent workability and having a tensile strength of 780 MPa or more; and methods of manufacturing the steel sheets.
  • TRIP-aided bainitic ferrite (TBF) steel sheet having a bainitic ferrite matrix and containing retained austenite (hereinafter also referred to as "retained ⁇ ") is known as one of TRIP-aided steel sheets (Patent Literature (PTL) 1 to 4).
  • the TBF steel sheet can have a high strength due to the hard bainitic ferrite and an elongation (EL) and a stretch flangeability ( ⁇ ) both at satisfactory levels due to the fine retained austenite present at the boundary of the bainitic ferrite.
  • Techniques disclosed in PTL 5 and 6 are known as techniques for increasing the elongation and stretch flangeability to thereby improve the workability.
  • the technique disclosed in PTL 5 allows the steel sheet to have a higher strength by the action of the martensitic microstructure and to have better workability by forming retained austenite in a predetermined amount.
  • the technique disclosed in PTL 6 allows the steel sheet to have a higher strength by the action of a lower bainitic microstructure and/or a martensitic microstructure and to have better workability by forming retained austenite and tempered martensite in predetermined amounts.
  • the techniques in PTL 5 and 6 control polygonal ferrite area percentage to be 10% or less so as to ensure a tensile strength of 980 MPa or more.
  • Dual-phase (DP) steel sheets and TRIP-aided steel sheets are known as steel sheets having a strength and workability both at certain levels.
  • the DP steel sheets have a metal structure including ferrite and martensite.
  • the TRIP-aided steel sheets utilize transformation induced plasticity (TRIP) of retained austenite (retained y).
  • a steel sheet disclosed in PTL 7 is known as one of such TRIP-aided steel sheets having a strength and workability both at certain levels.
  • PTL 7 discloses a technique of allowing the steel sheet to have, as a metal structure, a composite microstructure including martensite and retained austenite both present in ferrite and to thereby have a higher strength and better workability (particularly, a higher elongation).
  • PTL 8 discloses a technique of allowing a TRIP-aided steel sheet to have a better balance between a strength (TS) and an elongation (EL) (specifically, the product of TS and EL) to thereby have better stamping performance.
  • the technique disclosed in this literature allows the steel sheet to have a metal structure including ferrite, retained austenite, and at least one of bainite and martensite so as to have better stamping performance.
  • JP 2008 308717 discloses a high-strength steel sheet having excellent workability in which ductility and hole expansibility are consistent, comprising 0.08 to 0.3% C, 0.25 to 1.1% Si, greater than 2 to 3.5% Mn, 0.028% P or less, 0.01% S or less, 0.03 to 1% Al, 0.01% N or less and optionally, 1% Cu or less, 1% Ni or less, 1% Cr or less, 0.5% Mo or less and 0.005% B or less.
  • the steel sheet has a structure containing ferrite, a low temperature transformation phase and retained austenite, wherein the volume ratio of ferrite (VF) is 0.1 to 0.80, and retained austenite (VA), 0.30 or less. Further, the grain size of carbides present in the low temperature transformed phase is 500 nm or less and the following inequalities are satisfied: 0.2 ⁇ (Si+Al)/(Mn+Ni) ⁇ 0.80 and 0.15 ⁇ C/(1-VF) ⁇ 0.50.
  • JP S61 217529 discloses a steel sheet having superior tensile strength, ductility and secondary workability, by heat treating a steel sheet under specific conditions.
  • the steel sheet comprises by weight, 0.12-0.70% C, 0.4-1.8% Si, 0.2-2.5% Mn, 0.01-0.07% Al and less than 0.02% N.
  • the method comprises: heating to ⁇ Ac 3 at a rate of 1-100°C/sec; holding and annealing for ⁇ 3 min, preferably ⁇ 40sec; cooling the sheet to 350-500°C at a rate of 1-200°C/sec; aging at the temperature for 30 sec-10 min; cooling to 150-250°C at a rate of ⁇ 50°C/sec; cooling to normal temperature.
  • the steel sheet has a structure containing at least 1-50% ferrite and ⁇ 10% residual austenite.
  • steel sheets for use typically in pillars and members should be subjected to bulging or drawing under more severe conditions, because more and more demands have recently been made on steel sheet workability.
  • the steel sheets desirably have further better workability, particularly a further higher elongation.
  • local formability such as stretch flangeability ( ⁇ ) and bendability (R)
  • ⁇ stretch flangeability
  • R bendability
  • TRIP-aided steel sheets desirably have better local formability, such as stretch flangeability ( ⁇ ) and bendability (R), without deterioration in strength and elongation.
  • the aforementioned TRIP-aided steel sheets are disadvantageously poor in local formability, such as stretch flangeability and bendability, because retained austenite transforms into extremely hard martensite during working.
  • the present invention has been made under these circumstances, and an object thereof is to provide high-strength steel sheets excellent in workability, which have a tensile strength of 780 MPa or more and have an elongation and local formability both at higher levels. Another object of the present invention is to provide methods of manufacturing the steel sheets.
  • the present invention has achieved the objects and provides a high-strength steel sheet.
  • This high-strength steel sheet contains: C in a content of from 0.10% to 0.3%; Si in a content of from 1.2% to 3.0%; Mn in a content of from 1.5% to 3%; Al in a content of from 0.005% to 3%; P in a content controlled to 0.1% or less; and S in a content controlled to 0.05% or less, optionally further comprising: Cr in a content of from greater than 0% to 1%; and/or Mo in a content of from greater than 0% to 1%, optionally further comprising: at least one element selected from the group consisting of: Ti in a content of from greater than 0% to 0.15%; Nb in a content of from greater than 0% to 0.15%; and V in a content of from greater than 0% to 0.15%, optionally further comprising: Cu in a content of from greater than 0% to 1%; and/or Ni in a content of from greater than
  • the steel sheet has a metal structure including bainite, polygonal ferrite, retained austenite, and tempered martensite.
  • the bainite has a composite microstructure including: a high-temperature-formed bainite having an average distance between adjacent regions of retained austenite and/or carbide of 1 ⁇ m or more; and a low-temperature-formed bainite having an average distance between adjacent regions of retained austenite and/or carbide of less than 1 ⁇ m;
  • the high-temperature-formed bainite is present in an area percentage "a" of from 10% to 80% of the entire metal structure;
  • the low-temperature-formed bainite and the tempered martensite are present in a total area percentage "b" of from 10% to 80% of the entire metal structure;
  • the polygonal ferrite is present in an area percentage "c" of from 10% to 50% of the entire metal structure.
  • the retained austenite is present in a volume percentage of 5% or more of
  • martensite-austenite constituents including both as-quenched martensite and retained austenite are found upon observation at a cross-section of the metal structure with an optical microscope
  • martensite-austenite constituents each having an equivalent circle diameter "d" of greater than 7 ⁇ m are present in a number percentage of from 0% to less than 15% of a total number of the entire martensite-austenite constituents at the observed cross-section.
  • Grains of the polygonal ferrite preferably have an average equivalent circle diameter D of from greater than 0 ⁇ m to 10 ⁇ m.
  • the present invention further includes a high-strength hot-dip galvanized steel sheet including the high-strength steel sheet and, on a surface thereof, a hot-dip galvanized layer; and a high-strength hot-dip galvannealed steel sheet including the high-strength steel sheet and, on a surface thereof, a hot-dip galvannealed layer.
  • the high-strength steel sheet according to the present invention may be manufactured by a method including the steps of heating a steel sheet to a temperature range of from [ ⁇ (Ac 1 point+Ac 3 point)/2 ⁇ +20°C] to [Ac 3 point+20°C]; holding the steel sheet in the temperature range for 50 seconds or longer; cooling the steel sheet down to an arbitrary temperature T at an average cooling rate of 2°C or above per second, the temperature T falling within a range specified by Expression (1); holding the steel sheet in the temperature range (temperature range T1) specified by Expression (1) for 10 to 100 seconds; and holding the steel sheet in a temperature range (temperature range T2) specified by Expression (2) for 200 seconds or longer, in this order, Expressions (1) and (2) expressed as follows: 400 ° C ⁇ T 1 ° C ⁇ 540 ° C 200 ° C ⁇ T 2 ° C ⁇ 400 ° C
  • X and/or Y refers to that at least one of X and Y is contained.
  • the present invention can provide a high-strength steel sheet, which has elongation and local formability at satisfactory levels and exhibits superior workability even having a high strength of 780 MPa or more.
  • This is achieved by forming, particularly as bainite, both two different bainitic microstructures, i.e., a high-temperature-formed bainite and a low-temperature-formed bainite, and forming polygonal ferrite in a predetermined amount.
  • the high-temperature-formed bainite and the low-temperature-formed bainite differ from each other in existence form of retained austenite and carbides.
  • the high-temperature-formed bainite is formed at a high temperature in the range of from 400°C to 540°C; whereas the low-temperature-formed bainite is formed at a low temperature in the range of from 200°C to lower than 400°C.
  • the present invention can also provide a method of manufacturing the high-strength steel sheet having a strength and workability both at high levels.
  • a high-strength steel sheet according to the present invention will be illustrated.
  • the present inventors have made intensive investigations to improve the workability, particularly the elongation, and the local formability of a high-strength steel sheet having a tensile strength of 780 MPa or more. As a result, they have obtained findings (1) to (5) as follows:
  • the high-strength steel sheet according to the present invention has a metal structure as a mixed microstructure including bainite, polygonal ferrite, retained austenite, and tempered martensite.
  • bainite that features prominently in the present invention will be illustrated.
  • the "bainite” also includes bainitic ferrite.
  • the bainite microstructure includes precipitated carbides, whereas the bainitic ferrite microstructure does not include precipitated carbides.
  • a feature of the high-strength steel sheet according to the present invention is that the bainite has a composite microstructure including a high-temperature-formed bainite and a low-temperature-formed bainite having a higher strength than that of the high-temperature-formed bainite.
  • the high-temperature-formed bainite contributes to a higher steel sheet elongation, whereas the low-temperature-formed bainite contributes to better steel sheet local formability.
  • the steel sheet, as including the two different bainitic microstructures can have a higher elongation while surely having satisfactory local formability and exhibit better workability in general. This is probably because the combination of the bainitic microstructures differing in strength level causes nonuniform deformation and thereby increases the work hardenability.
  • high-temperature-formed bainite refers to a bainitic microstructure that is formed in the temperature range T1 of from 400°C to 540°C during a cooling process after heating to a temperature in the range of from [ ⁇ (Ac 1 point+Ac 3 point)/2 ⁇ +20°C] to [Ac 3 point+20°C] (two-phase temperature range).
  • the high-temperature-formed bainite may be defined as a microstructure having an average distance of the retained austenite or the like of 1 ⁇ m or more as determined by subjecting a cross section of the steel sheet to Nital etching and observing the etched cross section with a scanning electron microscope (SEM).
  • low-temperature-formed bainite refers to a bainitic microstructure that is formed in the temperature range T2 of from 200°C to lower than 400°C during the cooling process after heating to the two-phase temperature range.
  • the low-temperature-formed bainite may be defined as a microstructure having an average distance of the retained austenite or the like of less than 1 ⁇ m as determined by subjecting a cross section of the steel sheet to Nital etching and observing the etched cross section with a scanning electron microscope (SEM).
  • the term "average distance of the retained austenite or the like” refers to an average of center-to-center distances between adjacent retained austenite grains, center-to-center distances between adjacent carbide grains, and center-to-center distances between a retained austenite grain and an adjacent carbide as measured upon SEM observation of the steel sheet cross-section.
  • the term "center-to-center distance” refers to a distance (spacing) between center positions, which center positions are determined on each retained austenite or each carbide upon measurement on most adjacent regions of retained austenite and/or carbide. Each center position is determined by determining a major axis and a minor axis of the retained austenite or carbide and defining a location where the major axis intersects the minor axis as the center position.
  • the center-to-center distance may be defined not as a distance between adjacent grains of retained austenite and/or carbide but a spacing between one row and another which rows are formed by the retained austenite and/or carbide grains lying in a row in the major axis direction. This spacing is also referred to as a lath-to-lath distance.
  • the tempered martensitic microstructure functions similarly to the low-temperature-formed bainite and contributes to better steel sheet local formability.
  • the low-temperature-formed bainite and the tempered martensite are undistinguishable from each other even under SEM observation, and they are herein also synthetically referred to as a "low-temperature-formed bainite or the like".
  • the present invention can provide a high-strength steel sheet being improved in workability in general. This is achieved by allowing the steel sheet to have a composite bainitic microstructure including the high-temperature-formed bainite and the low-temperature-formed bainite or the like. Specifically, the high-temperature-formed bainite is softer than the low-temperature-formed bainite or the like and helps the steel sheet to have a higher elongation (EL) to thereby have better workability.
  • EL elongation
  • the low-temperature-formed bainite or the like is present as small grains of carbides and retained austenite, thereby relieves the stress concentration upon deformation, and helps the steel sheet to have a higher stretch flangeability ( ⁇ ) and/or a higher bendability (R) and to have better local formability. This contributes to better workability.
  • the high-strength steel sheet according to the present invention includes both the high-temperature-formed bainite and the low-temperature-formed bainite or the like, thereby has better work hardenability and a higher elongation, and exhibits better workability.
  • bainitic microstructures are classified as the "high-temperature-formed bainite” and the “the low-temperature-formed bainite or the like” by the differences in the forming temperature range and in the average distance of the retained austenite or the like. This is because the bainitic microstructures are not clearly distinguishable from each other by a general academic structural definition.
  • lathy bainite and bainitic ferrite are classified as upper bainite and lower bainite by the transformation temperature.
  • these microstructures, as well as the martensitic microstructure are undistinguishable from one another by SEM observation in steels containing Si in a high content of 1.0% or more as in the present invention.
  • the bainitic microstructures are classified or distinguished not by the academic structural definition but by the differences in the forming temperature range and in the average distance of the retained austenite or the like.
  • the high-temperature-formed bainite and the low-temperature-formed bainite or the like may be distributed in any distribution pattern not limited.
  • both the high-temperature-formed bainite and the low-temperature-formed bainite or the like may be present in each prior austenite grain; or the high-temperature-formed bainite and the low-temperature-formed bainite or the like may be formed in different prior austenite grains from each other (in this case, the high-temperature-formed bainite is present in one prior austenite grain, and the low-temperature-formed bainite or the like is present in another prior austenite grain).
  • Figs. 2 schematically illustrate how the high-temperature-formed bainite and the low-temperature-formed bainite or the like are distributed.
  • the high-temperature-formed bainite is hatched, and the low-temperature-formed bainite or the like is finely dotted.
  • Fig. 2(a) illustrates the high-temperature-formed bainite and the low-temperature-formed bainite or the like both present in each prior austenite grain.
  • Fig. 2(b) illustrates the high-temperature-formed bainite present in one prior austenite grain, and the low-temperature-formed bainite or the like present in another.
  • the filled circles illustrated in Figs. 2 represent martensite-austenite constituents, which will be described later.
  • the high-temperature-formed bainite is present in an area percentage of "a" of the entire metal structure; and the low-temperature-formed bainite or the like (the low-temperature-formed bainite and tempered martensite) is present in a total area percentage of "b" of the entire metal structure.
  • the area percentages "a” and "b” should each fall within the range of from 10% to 80%. Not an area percentage of the low-temperature-formed bainite, but a total area percentage of the low-temperature-formed bainite and the tempered martensite is specified herein. This is because these microstructures are undistinguishable from each other by SEM observation.
  • the area percentage "a” is controlled to the range of from 10% to 80%.
  • the high-temperature-formed bainite if present in an excessively small amount, may fail to help the steel sheet to have a sufficient elongation and to exhibit better workability.
  • the area percentage "a” is 10% or more, preferably 15% or more, and more preferably 20% or more.
  • the high-temperature-formed bainite if present in an excessively large amount, may fail to exhibit sufficient effects of the combination with the low-temperature-formed bainite or the like.
  • the high-temperature-formed bainite is present in an area percentage "a" of 80% or less, preferably 70% or less, more preferably 60% or less, and furthermore preferably 50% or less.
  • the total area percentage "b” is controlled to the range of from 10% to 80%.
  • the low-temperature-formed bainite or the like if present in an excessively small amount, may fail to help the steel sheet to have sufficient local formability and fail to improve the workability.
  • the total area percentage "b” is 10% or more, preferably 15% or more, and more preferably 20% or more.
  • the low-temperature-formed bainite or the like if present in an excessively large amount, may fail to exhibit sufficient effects of the combination with the high-temperature-formed bainite.
  • the low-temperature-formed bainite or the like is present in an area percentage "b" of 80% or less, preferably 70% or less, more preferably 60% or less, and furthermore preferably 50% or less.
  • the area percentage "a” and the total area percentage “b” may be in any relationship with each other, as long as they fall within the above-specified ranges respectively. All embodiments where “a” is larger than “b”; where “a” is smaller than “b”; and where “a” equals “b” are included herein.
  • the ratio between the high-temperature-formed bainite and the low-temperature-formed bainite or the like may be determined according to properties required of the steel sheet. Specifically, to further improve the local formability (particularly, the stretch flangeability ( ⁇ )) among the steel sheet workability, the ratio of the high-temperature-formed bainite may be minimized, whereas the ratio of the low-temperature-formed bainite or the like may be maximized. To further improve the elongation among the steel sheet workability, the ratio of the high-temperature-formed bainite may be maximized, whereas the ratio of the low-temperature-formed bainite or the like may be minimized. To further increase the steel sheet strength, the ratio of the low-temperature-formed bainite or the like may be maximized, whereas the ratio of the high-temperature-formed bainite may be minimized.
  • the polygonal ferrite microstructure is softer than bainite, functionally increases the steel sheet elongation, and contributes to better workability.
  • the polygonal ferrite is controlled to be present in an area percentage of 10% or more, preferably 12% or more, and more preferably 15% or more, of the entire metal structure.
  • the polygonal ferrite if present in an excessively large amount, may adversely affect the steel sheet strength.
  • the polygonal ferrite is present in an area percentage of 50% or less, preferably 45% or less, and more preferably 40% or less, of the entire metal structure.
  • the polygonal ferrite grains preferably have an average equivalent circle diameter D of from greater than 0 ⁇ m to 10
  • the steel sheet can have a further higher elongation by having a small average equivalent circle diameter D of the polygonal ferrite grains and allowing the polygonal ferrite to be dispersed as smaller grains. While the detailed mechanism thereof remains unclear, this is probably because the polygonal ferrite refinement allows the polygonal ferrite to be dispersed more uniformly in the entire metal structure, thereby impedes nonuniform deformation, and contributes to a further higher elongation.
  • the high-strength steel sheet according to the present invention has a mixed metal structure including the bainite, polygonal ferrite, retained austenite, and tempered martensite.
  • increasing sizes of the polygonal ferrite grains may cause variation in sizes of respective microstructures.
  • nonuniform deformation may occur to cause local strain concentration, and this may probably impede improvements in workability (particularly, effective improvement in elongation due to polygonal ferrite formation).
  • the polygonal ferrite grains have an average equivalent circle diameter D of preferably 10 ⁇ m or less, more preferably 8 ⁇ m or less, furthermore preferably 5 ⁇ m or less, and particularly preferably 3 ⁇ m or less.
  • the area percentage and average equivalent circle diameter D of the polygonal ferrite can be measured through SEM observation.
  • the total (a+b+c) of the area percentages "a", "b", and “c” is 70% or more of the entire metal structure.
  • the area percentage “a” is an area percentage of the high-temperature-formed bainite.
  • the area percentage “b” is a total area percentage of the low-temperature-formed bainite or the like (the low-temperature-formed bainite and tempered martensite).
  • the area percentage “c” is an area percentage of the polygonal ferrite.
  • the total area percentage (a+b+c) is more preferably 75% or more and furthermore preferably 80% or more.
  • the upper limit of the total area percentage (a+b+c) may be determined in consideration of the retained austenite volume percentage as determined by a saturation magnetization measurement, and is typically 95%.
  • the retained austenite transforms to martensite upon deformation of the steel sheet due to stress as received, effectively accelerates the deformed area to be hardened, and prevents strain concentration. This improves the uniform deformability and allows the steel sheet to exhibit a high elongation. These effects are generally known as "TRIP effects”.
  • the steel sheet contains the retained austenite in a volume percentage of 5% or more of the entire metal structure as determined by a saturation magnetization measurement.
  • the retained austenite may be present in a volume percentage of preferably 8% or more, and more preferably 10% or more.
  • the retained austenite if present in an excessively large amount, may cause excessive formation and resulting coarsening of the martensite-austenite constituents and adversely affect the local formability (stretch flangeability and bendability).
  • the upper limit of the retained austenite volume percentage may be about 30% and preferably 25%.
  • the retained austenite is predominantly formed between laths in the metal structure, but may be present as blocks as part of the martensite-austenite constituents on lathy microstructure aggregates (e.g., blocks and/or packets) or at the prior austenite grain boundary.
  • the high-strength steel sheet according to the present invention has a metal structure including bainite, polygonal ferrite, retained austenite, and tempered martensite as described above.
  • the metal structure may include these microstructures exclusively, but may further include any of other microstructures within ranges not adversely affecting advantageous effects of the present invention.
  • the other microstructures are exemplified by (a) a martensite-austenite constituent including both as-quenched martensite and retained austenite in combination; and (b) pearlite.
  • the martensite-austenite constituent (MA constituent) is generally known as a composite phase of as-quenched martensite (fresh martensite) and retained austenite and is a microstructure formed so that part of a microstructure, which has been present as untransformed austenite before final cooling, transforms to martensite during the final cooling, and the remainder remains as austenite.
  • the resulting martensite-austenite constituent is a very hard microstructure because carbon is enriched to a high concentration in the microstructure during a heat treatment process (particularly, austemper process) and part of the microstructure transforms to a martensitic microstructure.
  • the martensite-austenite constituent may readily cause voids as a result of stress concentration thereon during deformation.
  • the martensite-austenite constituent if present in an excessively large amount, may cause the steel sheet to have insufficient stretch flangeability and/or bendability, resulting in insufficient local formability.
  • the martensite-austenite constituent if present in an excessively large amount, may readily cause the steel sheet to have an excessively high strength.
  • the martensite-austenite constituent more readily forms with an increasing retained austenite amount and an increasing Si content. The amount of the formed martensite-austenite constituent is preferably minimized.
  • the metal structure may include the martensite-austenite constituents in an area percentage of preferably 30% or less, more preferably 25% or less, and furthermore preferably 20% or less, of the entire metal structure upon observation of the metal structure with an optical microscope.
  • martensite-austenite constituents each having an equivalent circle diameter "d" of greater than 7 ⁇ m are present in a number percentage of from 0% to less than 15% of the total number of the entire martensite-austenite constituents.
  • Such coarse martensite-austenite constituents each having an equivalent circle diameter "d” of greater than 7 ⁇ m may adversely affect the local formability.
  • the number percentage of the martensite-austenite constituents each having an equivalent circle diameter "d" of greater than 7 ⁇ m is more preferably less than 10% and furthermore preferably less than 5% of the total number of the entire martensite-austenite constituents.
  • the number percentage of the martensite-austenite constituents each having an equivalent circle diameter "d" of greater than 7 ⁇ m may be determined under observation of a surface of a cross-section in parallel to the rolling direction with an optical microscope.
  • martensite-austenite constituents are recommended to be minimized in size because it has been experimentally demonstrated that martensite-austenite constituents, when having large grain sizes, may readily cause void formation.
  • the metal structure may include the pearlite in an area percentage of 20% or less of the entire metal structure upon SEM observation of the metal structure.
  • the pearlite if contained in an area percentage of greater than 20%, may cause the steel sheet to have an insufficient elongation, and this may impede improvements in workability.
  • the area percentage of the pearlite is more preferably 15% or less, furthermore preferably 10% or less, and particularly preferably 5% or less of the entire metal structure.
  • the metal structure may be determined or measured by procedures as follows.
  • the high-temperature-formed bainite, the low-temperature-formed bainite or the like may be identified by subjecting a cross-section parallel to the steel sheet rolling direction at a depth of one-fourth the sheet thickness to Nital etching, and observing the etched cross-section with a SEM at a magnification of about 3000 times.
  • the high-temperature-formed bainite and the low-temperature-formed bainite or the like are observed mainly as gray microstructures, in which the retained austenite or the like is dispersed in the grains.
  • the retained austenite or the like is observed as a white or pale gray microstructure.
  • the high-temperature-formed bainite and the low-temperature-formed bainite or the like are therefore observed as including the retained austenite and carbides, and the area percentage thereof is calculated as one including the retained austenite or the like.
  • the polygonal ferrite is observed as grains that do not include the retained austenite or the like.
  • the retained austenite or the like, if present, is observed as a white or pale gray microstructure as described above.
  • the pearlite is observed as a layered microstructure including both carbides and ferrite.
  • the carbide and the retained austenite are observed both as white or pale gray microstructures and are undistinguishable from each other.
  • the carbide e.g., cementite
  • the carbide grains when dispersed with a large spacing, may be considered to be formed in a high temperature range; and, when dispersed with a small spacing, may be considered to be formed in a low temperature range.
  • the retained austenite is generally formed between laths. Such laths have smaller sizes with a decreasing forming temperature of the microstructure.
  • the retained austenite microstructures when dispersed with a large spacing, may be considered to be formed in a high temperature range; and, when dispersed with a small spacing, may be considered to be formed in a low temperature range.
  • the present invention therefore focuses attention on retained austenite or the like, which is observed as a white or pale gray microstructure in an observation view field upon SEM observation of a Nital-etched cross section of the steel sheet. Assume that center-to-center distances of adjacent retained austenite or the like are measured, and the measured distances are averaged to give an average (average center-to-center distance).
  • a microstructure having the average (average distance) of 1 ⁇ m or more is defined herein as a high-temperature-formed bainite; whereas a microstructure having the average (average distance) of less than 1 ⁇ m is defined as a low-temperature-formed bainite or the like.
  • the retained austenite microstructure is unidentifiable by SEM observation, and the volume percentage thereof is determined by a saturation magnetization measurement.
  • the measured volume percentage can be read as an "area percentage” thereof without modification.
  • a detailed measurement principle of the saturation magnetization measurement may be found in " R&D KOBE STEEL ENGINEERING REPORTS, Vol. 52, Sample No. 3, 2002, p. 43-46 ".
  • the martensite-austenite constituents may be observed as white microstructures by subjecting a cross-section parallel to the steel sheet rolling direction at a depth one-fourth the sheet thickness to Nital etching, and observing the etched cross-section with an optical microscope at a magnification of about 1000 times.
  • the high-strength steel sheet according to the present invention satisfactorily contains C in a content of from 0.10% to 0.3%; Si in a content of from 1.2% to 3.0%; Mn in a content of from 1.5% to 3%; and Al in a content of from 0.005% to 3% and is controlled to contain P in a content of from greater than 0% to 0.1%; and S in a content of from greater than 0% to 0.05%.
  • Carbon (C) element essentially helps the steel sheet to have a higher strength and allows the formation of retained austenite. For this reason, the carbon content is controlled to 0.10% or more, preferably 0.13% or more, and more preferably 0.15% or more. However, carbon, if contained in excess, may cause the steel sheet to have insufficient weldability. To prevent this, the carbon content is controlled to 0.3% or less, preferably 0.25% or less, and more preferably 0.20% or less.
  • Silicon (Si) element acts as a solute strengthening element, extremely importantly contributes to a higher strength of the steel sheet, suppresses precipitation of carbides during holding in the after-mentioned temperature range T1 and in the temperature range T2 (during austemper process), and thereby allows effective formation of the retained austenite.
  • the Si content is controlled to 1.2% or more, and more preferably 1.3% or more.
  • Si if contained in excess, may impede reverse transformation to the austenite phase during heating/soaking in annealing, and this may cause the steel sheet to have an insufficient strength due to a large amount of residual polygonal ferrite.
  • the excess Si may also cause silicon scales on the steel sheet surface during hot rolling and cause the steel sheet to have poor surface quality.
  • the Si content is controlled to 3.0% or less, preferably 2.5% or less, and more preferably 2.0% or less.
  • Manganese (Mn) element is necessary for obtaining bainite and tempered martensite.
  • the Mn element also effectively functionally stabilizes austenite ( ⁇ ) to form the retained austenite.
  • the Mn content is controlled to 1.5% or more, preferably 1.8% or more, and more preferably 2.0% or more.
  • Mn if contained in excess, may remarkably impede the formation of the high-temperature-formed bainite. Such excess Mn may also cause deterioration in weldability and in workability due to segregation. To prevent this, the Mn content is controlled to 3% or less, preferably 2.8% or less, and more preferably 2.7% or less.
  • Aluminum (Al) element suppresses the precipitation of carbides during the austemper process and contributes to the formation of the retained austenite, as with Si.
  • the Al element also acts as a deoxidizer during a steel-making process.
  • the Al content is controlled to 0.005% or more, preferably 0.01% or more, and more preferably 0.03% or more.
  • Al if contained in excess, may cause the steel sheet to contain excessively large amounts of inclusions to thereby have insufficient ductility.
  • the Al content is controlled to 3% or less, preferably 1.5% or less, more preferably 1% or less, and furthermore preferably 0.5% or less.
  • Phosphorus (P) element is an impurity to be inevitably contained in the steel. Phosphorus, if contained in excess, may cause the steel sheet to have insufficient weldability. To prevent this, the phosphorus content is controlled to 0.1% or less, preferably 0.08% or less, and more preferably 0.05% or less. The phosphorus content is desirably minimized, but is industrially hardly reduced to 0%.
  • S Sulfur element is an impurity to be inevitably contained in the steel and adversely affects the steel sheet weldability as with phosphorus. Sulfur forms sulfide inclusions in the steel sheet, and large amounts of such sulfide inclusions may adversely affect the weldability.
  • the sulfur content is controlled to 0.05% or less, preferably 0.01% or less, and more preferably 0.005% or less. The sulfur content is desirably minimized, but is industrially hardly reduced to 0%.
  • the high-strength steel sheet according to the present invention has a chemical composition within the above-specified ranges, with the remainder consisting of iron and inevitable impurities other than P and S.
  • the inevitable impurities are exemplified by N (nitrogen), O (oxygen), and tramp elements (e.g., Pb, Bi, Sb, and Sn).
  • N nitrogen
  • O oxygen
  • tramp elements e.g., Pb, Bi, Sb, and Sn.
  • the content of nitrogen is preferably from greater than 0% to 0.01%
  • the content of oxygen is preferably from greater than 0% to 0.01%.
  • Nitrogen (N) element precipitates as nitrides in the steel sheet and contributes to the steel sheet strengthening.
  • nitrogen if contained in excess, may cause precipitation of nitrides in large amounts and cause the steel sheet to deteriorate in elongation, stretch flangeability, and bendability.
  • the nitrogen content is preferably 0.01% or less, more preferably 0.008% or less, and furthermore preferably 0.005% or less.
  • Oxygen (O) element if contained in excess, may cause deterioration in elongation, stretch flangeability, and bendability.
  • the oxygen content is preferably 0.01% or less, more preferably 0.005% or less, and furthermore preferably 0.003% or less.
  • the high-strength steel sheet according to the present invention may further contain at least one element selected typically from:
  • rare-earth element(s) refers to and includes lanthanoid elements (fifteen elements ranging from La to Lu) as well as Sc (scandium) and Y (yttrium). Of these elements, at least one element selected from the group consisting of La, Ce, and Y is preferred, of which La and/or Ce is more preferred.
  • the high-strength steel sheet according to the present invention has a tensile strength of 780 MPa or more, has local formability and an elongation at satisfactory levels, and exhibits excellent workability.
  • the high-strength steel sheet is advantageously used as materials for automotive structural parts.
  • the automotive structural parts are exemplified by bumping parts such as front and rear side members and crush boxes; as well as pillars and other reinforcing members (e.g., center-pillar reinforcements); roof rail reinforcing members; side sills, floor members, kick-up parts, and other body-constituting parts; shock absorbing parts such as bumper reinforcing members and door impact beams; and seat parts.
  • the high-strength steel sheet also has good warm workability and is advantageously usable as materials for warm forming.
  • warm working (warm forming) refers to forming or shaping in a temperature range of from about 50°C to about 500°C.
  • the high-strength steel sheet can be manufactured by a method including the steps of heating a steel sheet to a temperature range of from [ ⁇ (Ac 1 point+Ac 3 point)/2 ⁇ +20°C] to [Ac 3 point+20°C] (two-phase temperature range), the steel sheet having a chemical composition within the above specified ranges; holding the steel sheet in the temperature range for 50 seconds or longer; cooling the steel sheet down to an arbitrary temperature T at an average cooling rate of 2°C or above per second, the temperature T falling within the range specified by Expression (1); holding the steel sheet in the temperature range specified by Expression (1) for 10 to 100 seconds; and holding the steel sheet in a temperature range specified by Expression (2) for 200 seconds or longer, in this order.
  • Expressions (1) and (2) are expressed as follows: 400 ° C ⁇ T 1 ° C ⁇ 540 ° C 200 ° C ⁇ T 2 ° C ⁇ 400 ° C
  • a high-strength steel sheet is prepared by hot-rolling a slab according to a common procedure to give a hot-rolled steel sheet; and cold-rolling the hot-rolled steel sheet.
  • This is prepared as a high-strength steel sheet to be heated to the two-phase temperature range [temperature range of from [ ⁇ (Ac 1 point+Ac 3 point)/2 ⁇ +20°C] to [Ac 3 point+20°C]].
  • the hot rolling may be performed at a finishing mill delivery temperature of typically 800°C or higher and a coiling temperature of typically 700°C or lower.
  • the cold rolling may be performed to a cold rolling reduction in a range of typically from 10% to 70%.
  • the heating to a temperature falling within the ferrite-austenite two-phase temperature range allows the polygonal ferrite formation in a predetermined amount. Specifically, heating to an excessively high temperature may cause the steel sheet to have a temperature falling within the austenite single-phase range, thereby suppress the formation of polygonal ferrite, and cause the steel sheet to fail to have a higher elongation and better workability.
  • the heating temperature is controlled to [Ac 3 point+20°C] or lower, preferably [Ac 3 point+10°C] or lower, and more preferably lower than the Ac 3 point.
  • the steel sheet, when heated to a temperature equal to or higher than the Ac 3 point has a temperature falling within the austenite single-phase temperature range.
  • the heating temperature is controlled to [ ⁇ (Ac 1 point+Ac 3 point)/2 ⁇ +20°C] or higher, preferably [ ⁇ (Ac 1 point+Ac 3 point)/2 ⁇ +30°C] or higher, and more preferably [ ⁇ (Ac 1 point+Ac 3 point)/2 ⁇ +50°C] or higher.
  • Soaking in the two-phase temperature range may fail to heat the steel sheet uniformly, and this may impede the formation of the retained austenite and cause the steel sheet to have an elongation and local formability at insufficient levels and to exhibit unsatisfactory workability.
  • the soaking time is controlled to 50 seconds or longer and preferably 100 seconds or longer.
  • soaking if performed for an excessively long time, may cause the austenite grains to have larger sizes, sequentially cause the polygonal ferrite grains to have larger sizes, and readily cause the steel sheet to have an elongation and local formability at insufficient levels.
  • the soaking time is preferably 500 seconds or shorter and more preferably 450 seconds or shorter.
  • the heating of the cold-rolled steel sheet to the two-phase temperature range may be performed at an average heating rate of typically 1°C or above per second.
  • the Ac 1 point and the Ac 3 point can be calculated from following Expressions (a) and (b), which are described in " Leslie Tekko Zairyo Kagaku (in Japanese; translation of "The Physical Metallurgy of Steels” by Leslie)", Maruzen Co., Ltd., May 31,1985, p. 273 .
  • a numerical value in brackets represents the content (in mass percent) of each element. When an element is not contained in the steel sheet, calculation can be performed while making the content of the element be 0% (in mass percent).
  • Cooling in this temperature range if performed at an average cooling rate of lower than 2°C per second, may cause pearlite transformation to form pearlite in excess, and this may cause the steel sheet to have an insufficient elongation and to exhibit unsatisfactory workability.
  • the average cooling rate in this temperature range is preferably 5°C or above per second, and more preferably 10°C or above per second.
  • the upper limit of the average cooling rate in this temperature range is not critical, but may be typically about 100°C per second, because cooling, if performed at an excessively high average cooling rate, may impede the temperature control.
  • the austemper process of holding the steel sheet in the temperature range T2 for a predetermined time allows untransformed austenite to transform to the low-temperature-formed bainite or martensite and allows carbon to be enriched in austenite to form retained austenite. This gives a metal structure as specified in the present invention.
  • the combination of holding in the temperature range T1 and holding in the temperature range T2 also effectively suppresses the formation of martensite-austenite constituents.
  • a mechanism of this is probably as follows. Si and Al, when added, generally suppress carbide precipitation. This allows the presence of free carbon in the steel and induces a phenomenon during the austemper process, in which the free carbon is enriched in the untransformed austenite along with bainitic transformation. The enrichment of carbon in the untransformed austenite allows the retained austenite to be formed in a large amount.
  • the steel sheet is held in the temperature range T2 after being held in the temperature range T1 according to the present invention.
  • This allows carbon to be enriched in a larger content in the untransformed austenite.
  • This also allows the bainitic transformation to proceed more in a low temperature range than in a high temperature range and helps the martensite-austenite constituents to be smaller.
  • Lathy microstructures become smaller in size in the case of holding the steel sheet in the temperature range T2 than in the case of holding the steel sheet in the temperature range T1. This helps the martensite-austenite constituents themselves, even when present, to be finely divided.
  • the high-temperature-formed bainite has already been formed at the time when holding of the steel sheet in the temperature range T2 starts, because the steel sheet is held in the temperature range T2 after being held in the temperature range T1 for a predetermined time.
  • the high-temperature-formed bainite triggers and accelerates the transformation of the low-temperature-formed bainite in the temperature range T2, and this advantageously effectively shortens the austemper process time.
  • the size reduction of the lathy microstructures and resulting size reduction of the martensite-austenite constituents can also be achieved by a simple austemper process of holding a steel sheet at a low temperature.
  • the steel sheet is not held in the temperature range T1 but cooled from the two-phase temperature range directly to an arbitrary temperature falling within the range specified by Expression (2) and held only in the temperature range T2 specified by Expression (2).
  • the high-temperature-formed bainite is little formed, and the lathy microstructure in the matrix have a large dislocation density because the steel sheet is not held in the temperature range T1.
  • the resulting steel sheet has an elongation and local formability at insufficient levels and exhibits poor workability.
  • the lower limit of the temperature range T1 is controlled to 400°C and preferably 420°C.
  • the steel sheet is held in the temperature range T1 for a holding time of from 10 to 100 seconds.
  • the holding if performed for a holding time of longer than 100 seconds, may cause excessive formation of the high-temperature-formed bainite, and this may impede the formation of a sufficient amount of the low-temperature-formed bainite or the like even when the steel sheet is subsequently held in the temperature range T2 for a predetermined time in a manner described later. This may cause the steel sheet to fail to have a strength and workability both at satisfactory levels.
  • Such long-term holding in the temperature range T1 may also cause excessive enrichment of carbon in the austenite.
  • the holding time is controlled to 100 seconds or shorter, preferably 90 seconds or shorter, and more preferably 80 seconds or shorter.
  • holding in the temperature range T1 if performed for an excessively short time, may cause the high-temperature-formed bainite to be formed in a smaller amount. This may cause the steel sheet to have an insufficient elongation and to exhibit unsatisfactory workability.
  • the holding time in the temperature range T1 is controlled to 10 seconds or longer, preferably 15 seconds or longer, more preferably 20 seconds or longer, and furthermore preferably 30 seconds or longer.
  • holding time in the temperature range T1 refers to a time (duration) from the time point when the steel sheet surface temperature reaches the upper limit temperature of the temperature range T1 to the time point when the steel sheet surface temperature reaches the lower limit temperature of the temperature range T1.
  • the term refers to a time from the time point when steel sheet surface temperature reaches 540°C to the time point when it reaches 400°C.
  • any of heat patterns (i), (ii), and (iii) illustrated in Fig. 3 may be employed.
  • the heat pattern (i) in Fig. 3 is an embodiment, in which the steel sheet is rapidly cooled from the two-phase temperature range down to an arbitrary temperature T falling within the range specified by Expression (1) and is then isothermally held at the temperature T for a predetermined time. After being isothermally held, the steel sheet is cooled down to an arbitrary temperature falling within the range specified by Expression (2).
  • the heat pattern (i) in Fig. 3 illustrates a one-stage isothermal holding. Isothermal holding, however, is not limited to this, and an isothermal holding in two or more stages at different holding temperatures may also be performed, as long as the steel sheet surface temperature fall within the temperature range T1.
  • the heat pattern (ii) in Fig. 3 is an embodiment, in which the steel sheet is rapidly cooled from the two-phase temperature range down to an arbitrary temperature T falling within the range specified by Expression (1), then cooled in the temperature range T1 at a changed cooling rate over a predetermined time, and cooled down to an arbitrary temperature falling within the range specified by Expression (2) at a further changed cooling rate.
  • the heat pattern (ii) in Fig. 3 illustrates an embodiment where the steel sheet is cooled in the temperature range T1 over a predetermined time.
  • the heat pattern in the present invention is not limited to this, and the process may further include the step of heating over a predetermined time or may further include the step of suitably repeating cooling and heating, as long as the steel sheet surface temperature fall within the temperature range T1.
  • a multistage cooling in which cooling is performed in two or more stages at different cooling rates, may be employed.
  • a one-stage heating or a multistage heating (two or more stages) may also be performed (not shown).
  • the heat pattern (iii) in Fig. 3 is an embodiment, in which the steel sheet is rapidly cooled from the two-phase temperature range down to an arbitrary temperature T falling within the range specified by Expression (1) and is then slowly cooled down to an arbitrary temperature falling within the range specified by Expression (2) at a changed but constant cooling rate. Even such slow cooling is accepted herein, as long as the residence time of the steel sheet within the temperature range T1 be from 10 to 100 seconds.
  • the present invention is not limited to the heat patterns (i) to (iii) illustrated in Fig. 3 , but can also employ any of other heat patterns than above, as long as the conditions specified herein be satisfied.
  • the temperature range T2 herein specified by Expression (2) is specifically from 200°C to lower than 400°C. Holding the steel sheet in the temperature range for a predetermined time allows the untransformed austenite, which has not transformed in the temperature range T1, to transform to the low-temperature-formed bainite, or martensite. The holding performed for such a sufficient holding time allows the bainitic transformation to proceed, gives the retained austenite finally, and allows the martensite-austenite constituents to be finer.
  • the martensite is present as as-quenched martensite immediately after transformation, but is tempered during holding in the temperature range T2, and remains as tempered martensite.
  • the tempered martensite has properties similar to those of the low-temperature-formed bainite that is formed in such a temperature range where martensitic transformation occurs. However, holding at a temperature of 400°C or higher may cause the formation of coarse martensite-austenite constituents, and this may cause the steel sheet to have an elongation and local formability at insufficient levels and to fail to exhibit better workability.
  • the temperature range T2 is set to lower than 400°C, preferably 390°C or lower, and more preferably 380°C or lower.
  • the lower limit of the temperature range T2 is controlled to 200°C, preferably 250°C, and more preferably 280°C.
  • the steel sheet is held in the temperature range T2 specified by Expression (2) for a time of 200 seconds or longer.
  • the holding if performed for a time of shorter than 200 seconds, may cause the low-temperature-formed bainite or the like to be formed in an insufficient amount, and this may invite a low carbon content in the austenite, an insufficient amount of the retained austenite, and a larger amount of the as-quenched martensite.
  • the resulting steel sheet may thereby have an excessively high strength and have an insufficient elongation and poor local formability.
  • the steel sheet may fail to have a higher elongation due to the low carton content in the austenite and the insufficient amount of the retained austenite.
  • the steel sheet may fail to have better local formability because the martensite-austenite constituents formed in the temperature range T1 are not refined (not become smaller).
  • the holding time is controlled to 200 seconds or longer, preferably 250 seconds or longer, and more preferably 300 seconds or longer.
  • the upper limit of the holding time is not critical, holding, if performed for an excessively long time, may induce poor productivity and cause the enriched carbon to precipitate as carbides. This may impede the formation of retained austenite and cause the steel sheet to have a low elongation and poor workability.
  • the upper limit of the holding time may be set typically to 1800 seconds.
  • holding time in the temperature range T2 refers to a time (duration) from the time point at which the steel sheet surface temperature reaches the upper limit temperature of the temperature range T2 to the time point at which the steel sheet surface temperature reaches the lower limit temperature of the temperature range T2.
  • the term refers to a time from the time point at which the surface temperature reaches lower than 400°C to the time point at which it reaches 200°C.
  • a way to hold the steel sheet in the temperature range T2 is not limited, as long as the residence time of the steel sheet in the temperature range T2 be 200 seconds or longer.
  • the steel sheet may be isothermally held (held at a constant temperature), or may be cooled and/or heated within the temperature range T2.
  • a multistage holding at different holding temperatures may also be employed.
  • the steel sheet After being held in the temperature range T2 for a predetermined time, the steel sheet is cooled down to room temperature and yields a high-strength steel sheet according to the present invention.
  • the high-strength steel sheet may bear a hot-dip galvanized layer or a hot-dip galvannealed layer on a surface thereof.
  • the hot-dip galvanized layer or the hot-dip galvannealed layer may be formed under any conditions not limited, such as known conditions.
  • the hot-dip galvanized layer may be formed at a plating bath temperature of preferably from 400°C to 500°C and more preferably from 440°C to 470°C.
  • the plating bath may have any chemical composition not limited, and a known hot-dip galvanization bath can be employed.
  • a hot-dip galvannealed steel sheet may be manufactured by subjecting a hot-dip galvanized steel sheet bearing a hot-dip galvanized layer to an alloying treatment according to a common procedure.
  • the alloying treatment may be performed typically by holding the hot-dip galvanized steel sheet at a temperature of typically from about 450°C to about 600°C (particularly preferably from about 480°C to about 570°C) for a duration of from about 5 to about 30 seconds (particularly preferably from about 10 to about 25 seconds).
  • the alloying treatment may be performed typically with a heating furnace, a direct fire, or an infrared heating furnace.
  • the heating process is also not limited and is exemplified by gas heating, induction heater heating (heating with induction heating equipment), or another customary process.
  • the technique according to the present invention is advantageously applied particularly to thin steel sheets having a thickness of 3 mm or less.
  • the high-strength steel sheet according to the present invention has been described above.
  • Experimental slabs were prepared by subjecting steels having chemical compositions given in Tables 1 and 2 below (with the remainder consisting of iron and inevitable impurities other than P, S, N, and O) to vacuum ingot making.
  • the rare-earth elements (REM) in Tables 1 and 2 employed a misch metal containing about 50% of La and about 30% of Ce.
  • the prepared experimental slabs were sequentially subjected to hot rolling, cold rolling, and continuous annealing and yielded specimens. Specific conditions are as follows.
  • the experimental slabs were heated to 1250°C and held at this temperature for 30 minutes, hot-rolled to a rolling reduction of about 90% at a finishing mill delivery temperature of 920°C, and cooled from this temperature down to a coiling temperature of 500°C at an average cooling rate of 30°C, and coiled. After being coiled, the works were held at the coiling temperature (500°C) for 30 minutes, then cooled down to room temperature through furnace cooling, and yielded hot-rolled steel sheets having a thickness of 2.6 mm.
  • the resulting hot-rolled steel sheets were subjected to acid wash to remove surface scales, cold-rolled to a cold rolling reduction of 46%, and yielded cold-rolled steel sheets having a thickness of 1.4 mm.
  • the resulting cold-rolled steel sheets were heated to temperatures (°C) given in Tables 3 to 5, held for durations given in Tables 3 to 5 for soaking, cooled according to any of the following four patterns, subjected to continuous annealing, and yielded the specimens.
  • Cooling Pattern i corresponding to the pattern (i) in Fig. 3
  • a work was cooled down to a start temperature T (°C) given in Tables 3 to 5 at an average cooling rate (°C per second) given in Tables 3 to 5, held at this start temperature T (°C), then cooled down to a start temperature (°C) in the temperature range T2 given in Tables 3 to 5, and held at this start temperature.
  • Tables 3 to 5 indicate a residence time (second) in the temperature range T1; and a residence time (second) in the temperature range T2.
  • the tables also indicate a time (second) from the time point at which holding in the temperature range T1 was completed to the time point at which the work temperature reached the start temperature in the temperature range T2.
  • Cooling Pattern ii corresponding to the pattern (ii) in Fig. 3
  • a work was cooled down to a start temperature T (°C) given in Tables 3 to 5 at an average cooling rate (°C per second) given in Tables 3 to 5, subsequently cooled down to an end temperature (°C) given in Tables 3 to 5, cooled down to a start temperature (°C) in the temperature range T2 given in Tables 3 to 5, and held at this start temperature for a duration (second) given in Tables 3 to 5.
  • Tables 3 to 5 indicate a residence time (second) in the temperature range T1; and a residence time (second) in the temperature range T2.
  • the tables also indicate a time (second) from the time point at which holding in the temperature range T1 was completed to the time point at which the work temperature reached the start temperature in the temperature range T2.
  • Cooling Pattern iii corresponding to the pattern (iii) in Fig. 3
  • a work was cooled down to a start temperature T (°C) given in Tables 3 and 4 at an average cooling rate (°C per second) given in Tables 3 and 4, then cooled down to a start temperature (°C) in the temperature range T2 given in Tables 3 and 4, and held at this start temperature.
  • Tables 3 and 4 indicate a residence time (second) in the temperature range T1; and a residence time (second) in the temperature range T2.
  • Sample No. 8 in Table 3 was a sample, in which the work after the soaking was held at 420°C for 450 seconds and then cooled down to room temperature at an average cooling rate of 5°C per second without holding (without stopping).
  • the residence time in the temperature range T2 of this sample indicated in Table 3 refers to a time necessary to pass through the temperature range T2.
  • the residence time in the temperature range T1 of this sample indicated in Table 3 refers to a time necessary to pass through the temperature range T1.
  • Table 3 also indicates a residence time in the temperature range T1 (second) and a residence time in the temperature range T2 (second).
  • each of the asterisked values was out of the temperature range T1 or the temperature range T2 as specified in the present invention, but was indicated as a temperature in each column for convenience of explanation of the heat pattern.
  • the prepared specimens were subjected to metal structure observation and mechanical properties evaluation according to the following procedures.
  • area percentages of the high-temperature-formed bainite, the low-temperature-formed bainite or the like were calculated based on an observation with a scanning electron microscope (SEM); whereas a volume percentage of the retained austenite was determined by a saturation magnetization measurement.
  • a specimen cross section parallel to the rolling direction was surface-polished, further electropolished, etched with a Nital solution, and a position at a one-fourth depth the sheet thickness was observed in five view fields with a SEM at a 3000-fold magnification.
  • One view field to be observed had a size of about 50 ⁇ m long by about 50 ⁇ m wide.
  • Tables 6 to 8 indicate an area percentage "a” (%) of the high-temperature-formed bainite, a total area percentage “b” (%) of the low-temperature-formed bainite and tempered martensite, and an area percentage “c” (%) of the polygonal ferrite.
  • the tables also indicate a total area percentage (a+b+c) of the area percentage "a”, the total area percentage "b", and the area percentage "c”.
  • a volume percentage of the retained austenite was determined by a saturation magnetization measurement. Specifically, a saturation magnetization (I) of a specimen and a saturation magnetization (Is) of a standard specimen were measured, based on which a volume percentage (Vyr) of the retained austenite was determined according to an expression below.
  • the standard specimen had been subjected to a heat treatment at 400°C for 15 hours.
  • the saturation magnetization measurement was performed at room temperature with a DC magnetization B-H curve automatic recorder "Model BHS-40" supplied by Riken Denshi Co., Ltd. with a maximum applied magnetization of 5000 (Oe).
  • Equivalent circle diameters "d" of martensite-austenite constituents including both retained austenite and as-quenched martensite in combination were measured by polishing a surface of a specimen cross section parallel to the rolling direction, and observing the polished surface in five view fields with an optical microscope at a 1000-fold magnification. A number percentage of martensite-austenite constituents each having an equivalent circle diameter "d" of greater than 7 ⁇ m with respect to the total number of the entire martensite-austenite constituents in the observed cross-section was calculated. A specimen having a number percentage of less than 15% was evaluated as accepted ( ⁇ ), whereas a specimen having a number percentage of 15% or more was evaluated as rejected ( ⁇ ). The evaluation results are indicated in Tables 6 to 8.
  • the V-bending test was performed at an angle between a die and a punch of 90° and at a punch tip radius varying in units of 0.5 mm.
  • a minimum punch tip radius at which the test specimen can be bent without cracking was determined as the critical bending radius (R).
  • R critical bending radius
  • the measurement results are indicated in Tables 6 to 8. The presence/absence of cracking was observed with a magnifying glass, and a test specimen without the occurrence of hair crack was evaluated as suffering from no cracking.
  • the Erichsen value was measured through an Erichsen test according to JIS Z 2247. A test specimen of a size of 90 mm long by 90 mm wide by 1.4 mm thick was prepared by cutting from the specimen. The Erichsen test was performed with a device having a punch diameter of 20 mm. The measurement results are indicated in Tables 6 to 8. The Erichsen test enables evaluations of both the total elongation properties and the local ductility of the steel sheet.
  • the mechanical properties of the specimen were evaluated based on criteria for the elongation (EL), bore expansion ratio ( ⁇ ), critical bending radius (R), and Erichsen value varying depending on the tensile strength (TS). Specifically, required properties including EL, ⁇ , R, and Erichsen value vary depending on the steel sheet tensile strength TS. Accordingly, the mechanical properties were evaluated in accordance with the following criteria that vary depending on the TS level.
  • the high-strength steel sheet is designed to be applied to steel sheets having a tensile strength TS of from 780 MPa to less than 1370 MPa. Samples having a tensile strength TS of less than 780 MPa, or of 1370 MPa or more were ineligible herein even when being satisfactory in EL, ⁇ , R, and Erichsen value.
  • Tables 1 to 8 demonstrate as follows. Of Samples Nos. 1 to 70 indicated in Tables 6 to 8, Samples Nos. 4, 29, 31, 38, 55, 65, and 67 underwent cooling in Pattern i; Samples Nos. 7, 11, 14, and 33 underwent cooling in Pattern iii; Samples Nos. 8 and 15 underwent cooling in Pattern iv; and the other samples underwent cooling in Pattern ii.
  • Samples Nos. 8, 13, 15, 29, 31, 34, 37, 41, 46, 48, 52, 57 and 60 to 63 in Tables 6 to 8) were steel sheets not satisfying one or more of the conditions specified in the present invention. Details thereof are as follows.
  • Sample No. 8 in Table 6 underwent holding in the temperature range T1 for an excessively long time and subsequent cooling without holding in the temperature range T2.
  • This sample suffered from insufficient formation of the low-temperature-formed bainite or the like and heavy formation of coarse martensite-austenite constituents.
  • the sample thereby had a low bore expansion ratio ⁇ and failed to exhibit better workability.
  • Sample No. 13 in Table 6 underwent, after being heated to and held in the two-phase temperature range, cooling down to an arbitrary temperature T falling within the range specified by Expression (1) at an excessively low average cooling rate.
  • Sample No. 31 in Table 6 underwent holding in the temperature range T1 for an excessively short time. This sample suffered from formation of an excessively small amount of the high-temperature-formed bainite, thereby had a low elongation, and failed to exhibit better workability.
  • Sample No. 34 in Table 7 underwent holding in the temperature range T1 for an excessively long time without holding in the temperature range T2. This sample suffered from insufficient formation of the low-temperature-formed bainite or the like and heavy formation of coarse martensite-austenite constituents. The sample thereby had a low Erichsen value and poor local formability and failed to exhibit better workability.
  • Sample No. 37 in Table 7 underwent heating at an excessively high temperature.
  • Sample No. 46 in Table 7 underwent holding in the two-phase temperature range for an excessively short time and suffered from insufficient formation of the retained austenite. This sample had a low elongation and a low Erichsen value to exhibit poor local formability, and failed to exhibit better steel sheet workability.
  • Sample No. 48 in Table 7 underwent, after soaking, holding at a temperature higher than the temperature range T1 specified in the present invention, subsequently cooling down to the temperature range T2 without holding in the temperature range T1, and holding in the temperature range T2. This sample suffered from excessive formation of polygonal ferrite, insufficient formation of the high-temperature-formed bainite, thereby had a low elongation, and failed to exhibit better workability.
  • Sample No. 60 in Table 8 had an excessively low carbon content. This sample suffered from formation of an excessively small amount of the retained austenite, had a low elongation and a low Erichsen value, and failed to exhibit better workability.
  • Sample No. 61 in Table 8 had an excessively high Si content. This sample suffered from excessive formation of polygonal ferrite, suffered from insufficient formation of the high-temperature-formed bainite and the low-temperature-formed bainite or the like, and failed to have a strength surely at a desired level. Sample No. 62 in Table 8 had an excessively low Si content. This sample suffered from insufficient formation of the retained austenite, thereby had a low elongation, and failed to exhibit better workability. Sample No.
  • Fig. 4 demonstrates that steel sheets could have a higher elongation (EL) and exhibit further better workability by controlling the average equivalent circle diameter D of polygonal ferrite grains to 10 ⁇ m or less, even when having a tensile strength (TS) at an equivalent level.
  • EL elongation
  • TS tensile strength

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Claims (4)

  1. Tôle d'acier à haute résistance présentant une excellente aptitude au façonnage, comprenant :
    C dans une teneur comprise entre 0,10 % et 0,3 % ;
    Si dans une teneur comprise entre 1,2 % et 3,0 % ;
    Mn dans une teneur comprise entre 1,5 % et 3 % ;
    Al dans une teneur comprise entre 0,005 % et 3 % ;
    P dans une teneur contrôlée à 0,1 % ou moins ; et
    S dans une teneur contrôlée à 0,05 % ou moins,
    éventuellement comprenant en outre :
    Cr dans une teneur comprise entre plus de 0 % et 1 % ; et/ou
    Mo dans une teneur comprise entre plus de 0 % et 1 %,
    éventuellement comprenant en outre :
    au moins un élément choisi dans le groupe constitué par :
    Ti dans une teneur comprise entre plus de 0 % et 0,15 % ;
    Nb dans une teneur comprise entre plus de 0 % et 0,15 % ; et
    V dans une teneur comprise entre plus de 0 % et 0,15 %,
    éventuellement comprenant en outre :
    Cu dans une teneur comprise entre plus de 0 % et 1 % ; et/ou
    Ni dans une teneur comprise entre plus de 0 % et 1 %,
    éventuellement comprenant en outre :
    B dans une teneur comprise entre plus de 0 % et 0,005 %, et
    éventuellement comprenant en outre :
    au moins un élément choisi dans le groupe constitué par :
    Ca dans une teneur comprise entre plus de 0 % et 0,01 % ;
    Mg dans une teneur comprise entre plus de 0 % et 0,01 % ; et
    un élément des terres rares dans une teneur comprise entre plus de 0 % et 0,01 %, en pourcentage en masse, le reste étant constitué de fer et d'impuretés inévitables, caractérisée en ce que :
    la tôle d'acier a une structure métallique composée de bainite, de ferrite polygonale, d'austénite résiduelle et de martensite trempée ;
    (1) lorsque la structure métallique est observée avec un microscope électronique à balayage,
    (1a) la bainite ayant une microstructure composite comprenant :
    une bainite formée à haute température ayant une distance moyenne entre des régions adjacentes d'austénite et/ou de carbure résiduel de 1 µm ou plus ; et
    une bainite formée à basse température ayant une distance moyenne entre des régions adjacentes d'austénite et/ou de carbure résiduel inférieure à 1 µm,
    la bainite formée à haute température étant présente en un pourcentage de surface « a » compris entre 10 % et 80 % de la structure métallique totale ;
    la bainite formée à basse température et la martensite trempée étant présentes en un pourcentage de surface totale « b » compris entre 10 % et 80 % de la structure métallique totale ; et
    (1b) la ferrite polygonale étant présente en un pourcentage de surface « c » compris entre 10 % et 50 % de la structure métallique totale ; et
    (2) l'austénite résiduelle étant présente dans un pourcentage en volume de 5 % ou plus de la structure métallique totale, tel que déterminé par une mesure de magnétisation à saturation.
  2. Tôle d'acier à haute résistance selon la revendication 1, lorsque des constituants de martensite-austénite comprenant à la fois de la martensite trempée et de l'austénite résiduelle sont trouvés par observation au niveau d'une section transversale de la structure métallique avec un microscope optique, des constituants de martensite-austénite ayant chacun un diamètre circulaire équivalent « d » supérieur à 7 µm étant présents en un pourcentage numérique compris entre 0 % et moins de 15 % du total d'un nombre entier de constituants martensite-austénite au niveau de la section transversale observée.
  3. Tôle d'acier à haute résistance selon la revendication 1 ou 2, les grains de la ferrite polygonale ayant un diamètre de cercle équivalent moyen D compris entre plus de 0 µm et 10 µm.
  4. Tôle d'acier à haute résistance selon la revendication 1, comprenant une couche galvanisée à chaud ou une couche recuite après galvanisation à chaud sur sa surface.
EP12765664.3A 2011-03-31 2012-03-21 Tôle d'acier à haute résistance présentant une excellente aptitude au façonnage, et son procédé de fabrication Active EP2695961B1 (fr)

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JP2011197671A JP5685167B2 (ja) 2011-03-31 2011-09-09 加工性に優れた高強度鋼板およびその製造方法
JP2011197670A JP5685166B2 (ja) 2011-03-31 2011-09-09 加工性に優れた高強度鋼板およびその製造方法
PCT/JP2012/057210 WO2012133057A1 (fr) 2011-03-31 2012-03-21 Tôle d'acier à haute résistance présentant une excellente aptitude au façonnage, et son procédé de fabrication

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Families Citing this family (15)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP5728115B1 (ja) 2013-09-27 2015-06-03 株式会社神戸製鋼所 延性および低温靭性に優れた高強度鋼板、並びにその製造方法
JP6306481B2 (ja) * 2014-03-17 2018-04-04 株式会社神戸製鋼所 延性及び曲げ性に優れた高強度冷延鋼板および高強度溶融亜鉛めっき鋼板、並びにそれらの製造方法
WO2015151427A1 (fr) 2014-03-31 2015-10-08 Ｊｆｅスチール株式会社 Tôle d'acier laminée à froid à haute résistance et à haut coefficient d'élasticité et procédé de production s'y rapportant
JP6284813B2 (ja) * 2014-04-18 2018-02-28 株式会社神戸製鋼所 強冷間加工性と加工後の硬さに優れる熱延鋼板
MX2017005569A (es) * 2014-10-30 2017-06-23 Jfe Steel Corp Lamina de acero de alta resistencia y metodo para la fabricacion de la misma.
JP6282576B2 (ja) * 2014-11-21 2018-02-21 株式会社神戸製鋼所 高強度高延性鋼板
US10876181B2 (en) 2015-02-24 2020-12-29 Nippon Steel Corporation Cold-rolled steel sheet and method of manufacturing same
MX2018001280A (es) * 2015-07-31 2018-05-17 Nippon Steel & Sumitomo Metal Corp Lamina de acero de estructura tipo mixta de transformacion inducida por tension y metodo para fabricar la misma.
JPWO2017033222A1 (ja) * 2015-08-21 2018-07-12 新日鐵住金株式会社 鋼板
JP6237962B1 (ja) 2016-01-22 2017-11-29 Ｊｆｅスチール株式会社 高強度鋼板及びその製造方法
JP6762868B2 (ja) * 2016-03-31 2020-09-30 株式会社神戸製鋼所 高強度鋼板およびその製造方法
CN108652380B (zh) * 2018-08-16 2023-09-19 尚赫（天津）科技开发有限公司 一种带远红外功能的保温杯、远红外钢及其制备方法
KR102321295B1 (ko) * 2019-12-18 2021-11-03 주식회사 포스코 가공성이 우수한 고강도 강판 및 그 제조방법
CN113122770B (zh) * 2019-12-31 2022-06-28 宝山钢铁股份有限公司 低碳低成本超高强复相钢板/钢带及其制造方法
CN114959422A (zh) * 2022-06-06 2022-08-30 山东冀凯装备制造有限公司 一种高强度低合金贝氏体铸钢的制备方法

Family Cites Families (25)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS61217529A (ja) * 1985-03-22 1986-09-27 Nippon Steel Corp 延性のすぐれた高強度鋼板の製造方法
JPS61272321A (ja) * 1985-05-28 1986-12-02 Nippon Steel Corp 超高強度冷延鋼板の製造方法
JPH01230715A (ja) * 1987-06-26 1989-09-14 Nippon Steel Corp プレス成形性の優れた高強度冷延鋼板の製造方法
JP3527092B2 (ja) 1998-03-27 2004-05-17 新日本製鐵株式会社 加工性の良い高強度合金化溶融亜鉛めっき鋼板とその製造方法
JP4524850B2 (ja) * 2000-04-27 2010-08-18 Ｊｆｅスチール株式会社 延性および歪時効硬化特性に優れた高張力冷延鋼板および高張力冷延鋼板の製造方法
JP2001329340A (ja) * 2000-05-17 2001-11-27 Nippon Steel Corp 成形性の優れた高強度鋼板及びその製造方法
US7314532B2 (en) * 2003-03-26 2008-01-01 Kabushiki Kaisha Kobe Seiko Sho (Kobe Steel, Ltd.) High-strength forged parts having high reduction of area and method for producing same
JP4320198B2 (ja) * 2003-03-28 2009-08-26 日新製鋼株式会社 衝撃特性と形状凍結性に優れた高強度冷延鋼板の製造方法
JP4411221B2 (ja) 2004-01-28 2010-02-10 株式会社神戸製鋼所 伸び及び伸びフランジ性に優れた低降伏比高強度冷延鋼板およびめっき鋼板並びにその製造方法
JP2006089804A (ja) * 2004-09-24 2006-04-06 Nisshin Steel Co Ltd 縮管性に優れたインパネリインフォースメント用高強度電縫鋼管の製造方法
JP4445365B2 (ja) * 2004-10-06 2010-04-07 新日本製鐵株式会社 伸びと穴拡げ性に優れた高強度薄鋼板の製造方法
JP4716358B2 (ja) * 2005-03-30 2011-07-06 株式会社神戸製鋼所 強度と加工性のバランスに優れた高強度冷延鋼板およびめっき鋼板
JP5076434B2 (ja) 2005-10-06 2012-11-21 Ｊｆｅスチール株式会社 成形性および塗装後耐食性に優れた高強度冷延鋼板およびその製造方法
EP1975266B1 (fr) * 2005-12-28 2012-07-11 Kabushiki Kaisha Kobe Seiko Sho (Kobe Steel, Ltd.) Feuille d'acier ultra-resistante
CN100510143C (zh) * 2006-05-29 2009-07-08 株式会社神户制钢所 延伸凸缘性优异的高强度钢板
JP4974341B2 (ja) * 2006-06-05 2012-07-11 株式会社神戸製鋼所 成形性、スポット溶接性、および耐遅れ破壊性に優れた高強度複合組織鋼板
JP5030200B2 (ja) 2006-06-05 2012-09-19 株式会社神戸製鋼所 伸び、伸びフランジ性および溶接性に優れた高強度鋼板
KR101082680B1 (ko) * 2006-07-14 2011-11-15 가부시키가이샤 고베 세이코쇼 고강도 강판 및 그 제조 방법
JP5223360B2 (ja) * 2007-03-22 2013-06-26 Ｊｆｅスチール株式会社 成形性に優れた高強度溶融亜鉛めっき鋼板およびその製造方法
JP4910898B2 (ja) * 2007-06-13 2012-04-04 住友金属工業株式会社 高強度鋼板およびその製造方法
JP2010065272A (ja) 2008-09-10 2010-03-25 Jfe Steel Corp 高強度鋼板およびその製造方法
JP5418047B2 (ja) * 2008-09-10 2014-02-19 Ｊｆｅスチール株式会社 高強度鋼板およびその製造方法
JP5365112B2 (ja) * 2008-09-10 2013-12-11 Ｊｆｅスチール株式会社 高強度鋼板およびその製造方法
JP5883211B2 (ja) * 2010-01-29 2016-03-09 株式会社神戸製鋼所 加工性に優れた高強度冷延鋼板およびその製造方法
JP5488129B2 (ja) * 2010-03-31 2014-05-14 新日鐵住金株式会社 冷延鋼板およびその製造方法

Non-Patent Citations (1)

* Cited by examiner, † Cited by third party
Title
None *

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KR20130125829A (ko) 2013-11-19
KR20150050592A (ko) 2015-05-08
CN103459638B (zh) 2015-07-15
US20140044988A1 (en) 2014-02-13
CN103459638A (zh) 2013-12-18
EP2942416B1 (fr) 2017-06-07
EP2695961A4 (fr) 2014-12-17
CN104762565B (zh) 2017-04-12
KR101604963B1 (ko) 2016-03-18
US20160355920A1 (en) 2016-12-08
EP2695961A1 (fr) 2014-02-12
KR101574400B1 (ko) 2015-12-03
WO2012133057A1 (fr) 2012-10-04
CN104762565A (zh) 2015-07-08
EP2942416A1 (fr) 2015-11-11

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