EP2246456A1 - Hochfestes stahlblech und herstellungsverfahren dafür - Google Patents

Hochfestes stahlblech und herstellungsverfahren dafür Download PDF

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Publication number
EP2246456A1
EP2246456A1 EP09706046A EP09706046A EP2246456A1 EP 2246456 A1 EP2246456 A1 EP 2246456A1 EP 09706046 A EP09706046 A EP 09706046A EP 09706046 A EP09706046 A EP 09706046A EP 2246456 A1 EP2246456 A1 EP 2246456A1
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EP
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Prior art keywords
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steel sheet
martensite
high strength
temperature range
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EP09706046A
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English (en)
French (fr)
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EP2246456A4 (de
EP2246456B1 (de
EP2246456B9 (de
Inventor
Hiroshi Matsuda
Reiko Mizuno
Yoshimasa Funakawa
Yasushi Tanaka
Tatsuya Nakagaito
Saiji Matsuoka
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JFE Steel Corp
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JFE Steel Corp
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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
    • C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
    • C23C2/06—Zinc or cadmium or alloys based thereon
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0421—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the working steps
    • C21D8/0426—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0421—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the working steps
    • C21D8/0436—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0447—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
    • 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
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • 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
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
    • C23C2/0224—Two or more thermal pretreatments
    • 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
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C23C2/024—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by cleaning or etching
    • 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
    • C23C2/26—After-treatment
    • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
    • 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
    • C23C2/26—After-treatment
    • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
    • C23C2/29—Cooling or quenching
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/004—Dispersions; Precipitations
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0447—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the heat treatment
    • C21D8/0468—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the heat treatment between cold rolling steps

Definitions

  • the present invention relates to a high strength steel sheet that is used in industrial fields such as an automobile industry and an electrical industry, has good formability, and has a tensile strength of 900 MPa or higher and a method for manufacturing the same.
  • the high strength steel sheet of the present invention includes steel sheets whose surface is galvanized or galvannealed.
  • Patent Document 1 discloses a high strength steel sheet with a low yield ratio that is excellent in surface quality and bendability and has a tensile strength of 588 to 882 MPa and a method for manufacturing the steel sheet, by specifying the composition and the hot-rolling and annealing conditions.
  • Patent Document 2 discloses a high strength cold-rolled steel sheet with excellent bendability and a method for manufacturing the steel sheet, by specifying the hot-rolling, cold-rolling, and annealing conditions of steel having a certain composition.
  • Patent Document 3 discloses a steel sheet that is excellent in collision safety and formability and a method for manufacturing the steel sheet, by specifying the volume fraction and grain diameter of martensite and the mechanical properties.
  • Patent Document 4 discloses a high strength steel sheet, a high strength galvanized steel sheet, and a high strength galvannealed steel sheet that are excellent in stretch-flangeability and crashworthiness and a method for manufacturing the steel sheets, by specifying the composition and the volume fraction and grain diameter of martensite.
  • Patent Document 5 discloses a high strength steel sheet, a high strength galvanized steel sheet, and a high strength galvannealed steel sheet that are excellent in stretch-flangeability, shape fixability, and crashworthiness and a method for manufacturing the steel sheets, by specifying the composition, the grain diameter and microstructure of ferrite, and the volume fraction of martensite.
  • Patent Document 6 discloses a high strength steel sheet having excellent mechanical properties and a method for manufacturing the steel sheet, by specifying the composition, the amount of martensite, and the manufacturing method.
  • Patent Documents 7 and 8 each disclose a high strength galvanized steel sheet that is excellent in stretch-flangeability and bendability and a method and facility for manufacturing the steel sheet, by specifying the composition and the manufacturing conditions in a galvanizing line.
  • Patent Documents disclose steel sheets having a microstructure including a phase other than martensite as a hard second phase.
  • Patent Document 9 discloses a steel sheet that is excellent in fatigue properties, by employing martensite and/or bainite as a hard second phase and specifying the composition, the grain diameter, the hardness ratio, and the like.
  • Patent Document 10 discloses a steel sheet that is excellent in stretch-flangeability, by mainly employing bainite or pearlite as a second phase and specifying the composition and the hardness ratio.
  • Patent Document 11 discloses a high-strength and ductility galvanized steel sheet that is excellent in hole expandability and a method for manufacturing the steel sheet, by employing bainite and martensite as a hard second phase.
  • Patent Document 12 discloses a multiple-phase steel sheet that is excellent in fatigue properties by employing bainite and martensite as a hard second phase and specifying the fraction of constituent phases, the grain diameter, the hardness, and the mean free path of the entire hard phase.
  • Patent Document 13 discloses a high strength steel sheet that is excellent in ductility and hole expandability, by specifying the composition and the amount of retained austenite.
  • Patent Document 14 discloses a high strength multiple-phase cold-rolled steel sheet that is excellent in workability, by employing a steel sheet including bainite and retained austenite and/or martensite and specifying the composition and the fraction of phases.
  • Patent Document 15 discloses a high strength steel sheet that is excellent in workability and a method for manufacturing the steel sheet, by specifying the distribution state of the grains of a hard second phase in ferrite and the ratio of the grains of tempered martensite and bainite to the grains of ferrite.
  • Patent Document 16 discloses an ultra-high strength cold-rolled steel sheet that is excellent in delayed fracture resistance and has a tensile strength of 1180 MPa or higher and a method for manufacturing the steel sheet, by specifying the composition and the manufacturing process.
  • Patent Document 17 discloses an ultra-high strength cold-rolled steel sheet that is excellent in bendability and has a tensile strength of 980 MPa or higher and a method for manufacturing the steel sheet, by specifying the composition and the manufacturing method.
  • Patent Document 18 discloses an ultra-high strength thin steel sheet that has a tensile strength of 980 MPa or higher and whose hydrogen embrittlement is prevented by limiting the number of iron-based carbide grains in tempered martensite to a certain number and a method for manufacturing the steel sheet.
  • Patent Documents 1 to 7, 9 to 10, and 12 to 14 disclose the inventions regarding steel sheets having a tensile strength of lower than 900 MPa, and the workability often cannot be maintained if the strength is further increased.
  • Patent Document 1 describes that annealing is performed in a single phase region and the subsequent cooling is performed to 400°C at a cooling rate of 6 to 20 °C/s.
  • the adhesion of a coating needs to be taken into account and heating needs to be performed before coating because 400°C is lower than the temperature of a coating bath.
  • the galvanized steel sheet cannot be manufactured in a continuous galvanizing and galvannealing line having no heating equipment before the coating bath.
  • Patent Documents 7 and 8 since tempered martensite needs to be formed during the heat treatment in a galvanizing line, there is required equipment for reheating the steel sheet after the cooling to Ms temperature or lower.
  • Patent Document 11 bainite and martensite are employed as a hard second phase and the fraction is specified.
  • the characteristics significantly vary in the specified range, and the operating conditions need to be precisely controlled to suppress the variation.
  • Patent Document 15 since cooling is performed to Ms temperature or lower to form martensite before bainite transformation, equipment for reheating the steel sheet is required. Furthermore, the operating conditions need to be precisely controlled to achieve stable characteristics.
  • Patent Documents 16 and 17 the steel sheet needs to be maintained in a bainite-formation temperature range after annealing to obtain a microstructure mainly composed of bainite, which makes it difficult to achieve ductility.
  • the steel sheet In the case of a galvanized steel sheet, the steel sheet needs to be reheated to a temperature higher than the temperature of a coating bath.
  • Patent Document 18 only describes the improvement in hydrogen embrittlement of a steel sheet, and there is little consideration for workability although bendability is considered to some extent.
  • the ratio of a hard second phase to the entire microstructure needs to be increased to increase the strength of a steel sheet.
  • the ratio of a hard second phase is increased, the workability of a steel sheet is strongly affected by that of the hard second phase. The reason is as follows.
  • the ratio of the hard second phase is low, minimal workability is achieved by the deformation of ferrite itself that is a parent phase even if the workability of the hard second phase is insufficient.
  • the ratio of the hard second phase is high, the formability of a steel sheet is directly affected by the deformability of the hard second phase, not the deformation of ferrite. If the workability is insufficient, the formability is considerably degraded.
  • martensite is formed through water quenching by adjusting the fraction of ferrite and a hard second phase using a continuous annealing furnace that can perform water quenching. Subsequently, the temperature is increased and held to temper martensite, whereby the workability of the hard second phase is improved.
  • bainite When bainite is used, there is a problem in that the characteristics significantly vary due to the variation in a bainite-formation temperature range and the holding time.
  • martensite or retained austenite including bainite containing retained austenite
  • a mixed microstructure of martensite and bainite is considered to be used as a second phase microstructure to ensure both ductility and stretch-flangeability.
  • An object of the present invention is to provide a high strength steel sheet having a tensile strength of 900 MPa or higher that can minimize the formation of bainite, which easily causes a variation in characteristics such as strength and formability, and can have both high strength and good formability and to provide an advantageous method for manufacturing the high strength steel sheet.
  • TS ⁇ T. El The formability is evaluated using TS ⁇ T. El and a ⁇ value that represents stretch-flangeability.
  • TS ⁇ T. El ⁇ 14500 MPa ⁇ % and ⁇ ⁇ 15% are target characteristics.
  • the inventors of the present invention have studied about the formation process of martensite, in particular, the effect of the cooling conditions of a steel sheet on martensite.
  • a high strength steel sheet having both good formability and high strength with a tensile strength of 900 MPa or higher that are targeted in the present invention can be obtained by the following method.
  • martensite transformation is caused while at the same time the transformed martensite is tempered.
  • the ratio of the thus-formed autotempered martensite is controlled to a certain ratio and also the distribution state of iron-based carbide grains included in the autotempered martensite is suitably controlled, whereby such a high strength steel sheet can be obtained.
  • a high strength steel sheet having a tensile strength of 900 MPa or higher that can achieve high strength, good workability, and good ductility can be obtained by forming an appropriate amount of autotempered martensite in a steel sheet and suitably controlling the distribution state of carbide grains included in the autotempered martensite. Therefore, the present invention significantly contributes to the weight reduction of automobile bodies.
  • the present invention contributes to decreases in the number of steps and in the cost.
  • the ratio between ferrite and a hard phase described below is important and thus the area ratio of ferrite needs to be 5% or more and 80% or less. If the area ratio of ferrite is less than 5%, ductility is not ensured. If the area ratio of ferrite is more than 80%, the area ratio of the hard phase becomes insufficient and thus the strength becomes insufficient.
  • the area ratio of ferrite is preferably set in the range of 10% or more and 65% or less.
  • autotempered martensite is a microstructure obtained by simultaneously causing martensite transformation and the tempering of the martensite through autotempering treatment, and not so-called tempered martensite obtained through quenching and tempering treatments as in the related art.
  • the microstructure is not a uniformly tempered microstructure formed by completing martensite transformation through quenching and then performing tempering through a temperature increase as in typical quenching and tempering treatments, but is a microstructure including martensites in different tempered states obtained by performing martensite transformation and the tempering of the martensite in stages through the control of a cooling process in a temperature range of Ms temperature or lower.
  • This autotempered martensite is a hard phase for increasing strength. If the area ratio of autotempered martensite is less than 15%, sufficient strength cannot be achieved and work hardening of ferrite cannot be facilitated. Thus, the area ratio of autotempered martensite needs to be 15% or more and is preferably 30% or more.
  • the microstructure of a steel sheet is preferably composed of ferrite and autotempered martensite within the above-described range.
  • phases such phases are formed, other phases such as bainite, retained austenite, and as-quenched martensite are sometimes formed. These phases may be formed as long as some parameters are within the tolerable ranges described below. The tolerable ranges will now be described.
  • Bainite is a hard phase that contributes to an increase in strength, but the characteristics significantly vary in accordance with the formation temperature range and the variation in the quality of material is sometimes increased. Therefore, the area ratio of bainite in a steel microstructure is desirably as low as possible, but up to 10% of bainite is tolerable. The area ratio of bainite is preferably 5% or less.
  • Area ratio of retained austenite 5% or less (including 0%)
  • Retained austenite is transformed into hard martensite when processed, which decreases stretch-flangeability.
  • the area ratio of retained austenite in a steel microstructure is desirably as low as possible, but up to 5% of retained austenite is tolerable.
  • the area ratio of retained austenite is preferably 3% or less.
  • Area ratio of as-quenched martensite 40% or less (including 0%)
  • the area ratio of as-quenched martensite in a steel microstructure is desirably as low as possible, but up to 40% of as-quenched martensite is tolerable.
  • the area ratio of as-quenched martensite is preferably 30% or less.
  • as-quenched martensite can be differentiated from autotempered martensite in that carbides of as-quenched martensite are not observed with a scanning electron microscope (SEM) or a transmission electron microscope (TEM).
  • HV ⁇ 700 needs to be satisfied and HV ⁇ 630 is preferably satisfied.
  • Autotempered martensite is martensite subjected to the heat treatment (autotempering treatment) performed by the method of the present invention.
  • the mean hardness of autotempered martensite is HV ⁇ 700
  • the degree of autotempering treatment can be confirmed through the formation state (distribution state) of iron-based carbide grains in autotempered martensite.
  • the mean number of precipitated iron-based carbide grains each having a size of 5 nm or more and 0.5 ⁇ m or less is 5 ⁇ 10 4 or more per 1 mm 2 , it can be judged that desired autotempering treatment has been performed.
  • Iron-based carbide grains each having a size of less than 5 nm are removed from the target of judgment because such carbide grains do not affect the workability of autotempered martensite.
  • iron-based carbide grains each having a size of more than 0.5 ⁇ m are also removed from the target of judgment because such carbide grains may decrease the strength of autotempered martensite but hardly affect the workability. If the number of iron-based carbide grains is less than 5 ⁇ 10 4 per 1 mm 2 , it is judged that the autotempering treatment has been improperly performed because workability, particularly stretch-flangeability, is not improved.
  • the number of iron-based carbide grains is preferably 1 ⁇ 10 5 or more and 1 ⁇ 10 6 or less per 1 mm 2 , more preferably 4 ⁇ 10 5 or more and 1 ⁇ 10 6 or less per 1 mm 2 .
  • an iron-based carbide is mainly Fe 3 C, and ⁇ carbides and the like may be further contained.
  • Carbide grains can be identified by, for example, performing SEM-EDS (energy dispersive X-ray spectrometry), EPMA (electron probe microanalyzer), or FE-AES (field emission-Auger electron spectrometry) on samples whose section is polished.
  • SEM-EDS energy dispersive X-ray spectrometry
  • EPMA electron probe microanalyzer
  • FE-AES field emission-Auger electron spectrometry
  • the amount of autotempered martensite narrowed down by further limiting the size and number of iron-based carbide grains precipitated in the above-described autotempered martensite can be suitably set as follows.
  • Autotempered martensite in which the number of precipitated iron-based carbide grains each having a size of 0.1 ⁇ m or more and 0.5 ⁇ m or less is 5 ⁇ 10 2 or less per 1 mm 2 : the area ratio of the autotempered martensite to the entire autotempered martensite is 3% or more
  • the area ratio of autotempered martensite in which the number of precipitated iron-based carbide grains each having a size of 0.1 ⁇ m or more and 0.5 ⁇ m or less is 5 ⁇ 10 2 or less per 1 mm 2 to the entire autotempered martensite is preferably 3% or more.
  • the area ratio of such autotempered martensite to the entire autotempered martensite is preferably 40% or less, more preferably 30% or less.
  • the area ratio of autotempered martensite in which the number of precipitated iron-based carbide grains each having a size of 0.1 ⁇ m or more and 0.5 ⁇ m or less is 5 ⁇ 10 2 or less per 1 mm 2 to the entire autotempered martensite is 3% or more, the number of fine iron-based carbide grains is increased in autotempered martensite. Therefore, the mean number of precipitated iron-based carbide grains in the entire autotempered martensite is increased.
  • the mean number of precipitated iron-based carbide grains each having a size of 5 nm or more and 0.5 ⁇ m or less in autotempered martensite is preferably 1 ⁇ 10 5 or more and 5 ⁇ 10 6 or less per 1 mm 2 , more preferably 4 ⁇ 10 5 or more and 5 ⁇ 10 6 or less per 1 mm 2 .
  • the autotempered martensite microstructure includes a portion that contains a large number of iron-based carbide grains having a relatively large size and a portion that contains a small number of iron-based carbide grains having a relatively large size in a mixed manner.
  • the portion that contains a small number of iron-based carbide grains having a relatively large size is hard autotempered martensite because a large number of fine iron-based carbide grains are contained.
  • the portion that contains a large number of iron-based carbide grains having a relatively large size is soft autotempered martensite.
  • C is an essential element for increasing the strength of a steel sheet.
  • a C content of less than 0.1% causes difficulty in achieving both strength and workability such as ductility or stretch-flangeability of the steel sheet.
  • a C content of more than 0.3% causes a significant hardening of welds and heat-affected zones, thereby reducing weldability.
  • the C content is set in the range of 0.1% or more and 0.3% or less, preferably 0.12% or more and 0.23% or less.
  • Si is a useful element for solution hardening of ferrite, and the Si content is preferably 0.1% or more to ensure the ductility and the hardness of ferrite.
  • the excessive addition of Si causes the degradation of surface quality due to the occurrence of red scale and the like and the degradation of the adhesion of a coating.
  • the Si content is set to 2.0% or less, preferably 1.6% or less.
  • Mn 0.5% or more and 3.0% or less
  • Mn is an element that is effective in strengthening steel, stabilizes austenite, and is necessary for ensuring the area ratio of a hard phase. To achieve this, a Mn content of 0.5% or more is required. On the other hand, an excessive Mn content of more than 3.0% causes the degradation of castability or the like. Thus, the Mn content is set in the range of 0.5% or more and 3.0% or less, preferably 1.5% or more and 2.5% or less.
  • P causes embrittlement due to grain boundary segregation and degrades shock resistance, but a P content of up to 0.1% is tolerable. Furthermore, in the case where a steel sheet is galvannealed, a P content of more than 0.1% significantly reduces the rate of alloying. Thus, the P content is set to 0.1% or less, preferably 0.05% or less.
  • S is formed into MnS as an inclusion that causes the degradation of shock resistance and causes cracks along a flow of a metal in a weld zone.
  • the S content is preferably minimized.
  • a S content of up to 0.07% is tolerable in terms of manufacturing costs.
  • the S content is preferably 0.04% or less.
  • Al is an element that contributes to ferrite formation and a useful element for controlling the amount of the ferrite formation during manufacturing.
  • an excessive Al content degrades the quality of a slab during steelmaking.
  • the Al content is set to 1.0% or less, preferably 0.5% or less. Since an excessively low Al content sometimes makes it difficult to perform deoxidization, the Al content is preferably 0.01% or more.
  • N is an element that most degrades the anti-aging property of steel. Therefore, the N content is preferably minimized. A N content of more than 0.008% causes significant degradation of an anti-aging property. Thus, the N content is set to 0.008% or less, preferably 0.006% or less.
  • the steel sheet of the present invention can suitably contain the components described below in addition to the basic components described above.
  • Cr, V, and Mo have an effect of suppressing the formation of pearlite when a steel sheet is cooled from the annealing temperature and thus can be optionally added.
  • the effect is produced at a Cr content of 0.05% or more, a V content of 0.005% or more, or a Mo content of 0.005% or more.
  • an excessive Cr content of more than 5.0%, an excessive V content of more than 1.0%, or an excessive Mo content of more than 0.5% excessively increases the area ratio of a hard phase, thereby unnecessarily increasing the strength.
  • the Cr content is preferably set in the range of 0.005% or more and 5.0% or less
  • the V content is preferably set in the range of 0.005% or more and 1.0% or less
  • the Mo content is preferably set in the range of 0.005% or more and 0.5% or less.
  • At least one element selected from Ti, Nb, B, Ni, and Cu can be incorporated.
  • the reason for the limitation of the content ranges is as follows.
  • Ti 0.01% or more and 0.1% or less and Nb: 0.01% or more and 0.1% or less
  • Ti and Nb are useful for precipitation strengthening of steel and the effect is produced at a Ti content of 0.01% or more or a Nb content of 0.01% or more.
  • a Ti content of more than 0.1% or a Nb content of more than 0.1% degrades the workability and shape flexibility.
  • the Ti content and the Nb content are each preferably set in the range of 0.01% or more and 0.1% or less.
  • B has an effect of suppressing the formation and growth of ferrite from austenite grain boundaries and thus can be optionally added.
  • the effect is produced at a B content of 0.0003% or more.
  • a B content of more than 0.0050% decreases workability.
  • the B content is preferably set in the range of 0.0003% or more and 0.0050% or less.
  • the formation of BN is preferably suppressed to produce the above-described effect.
  • B is preferably added together with Ti.
  • Ni 0.05% or more and 2.0% or less
  • Cu 0.05% or more and 2.0% or less
  • Ni and Cu promote internal oxidation, thereby improving the adhesion of a coating.
  • the effect is produced at a Ni content of 0.05% or more or a Cu content of 0.05% or more.
  • a Ni content of more than 2.0% or a Cu content of more than 2.0% degrades the workability of a steel sheet.
  • Ni and Cu are useful elements for strengthening steel.
  • the Ni content and the Cu content are each preferably set in the range of 0.05% or more and 2.0% or less.
  • Ca and REM are useful elements for spheroidizing the shape of a sulfide and improving an adverse effect of the sulfide on stretch-flangeability.
  • the effect is produced at a Ca content of 0.001% or more or an REM content of 0.001% or more.
  • a Ca content of more than 0.005% or an REM content of more than 0.005% increases the number of inclusions or the like and causes, for example, surface defects and internal defects.
  • the Ca content and the REM content are each preferably set in the range of 0.001% or more and 0.005% or less.
  • components other than the components described above are Fe and incidental impurities.
  • a component other than the components described above may be contained to the extent that the advantages of the present invention are not impaired.
  • the composition of the steel sheet according to the present invention preferably satisfies M ⁇ 300°C that represents a relation between the composition and the area ratio of polygonal ferrite to perform stable production, that is, to suppress the variation in characteristics due to the variation in manufacturing conditions.
  • a galvanized layer or a galvannealed layer may be disposed on a surface of a steel sheet.
  • a slab prepared to have the above-described preferred composition is produced, hot-rolled, and then cold-rolled to obtain a cold-rolled steel sheet.
  • these processes are not particularly limited, and can be performed by typical methods.
  • a slab is heated to 1100°C or higher and 1300°C or lower and subjected to finish hot-rolling at a temperature of 870°C or higher and 950°C or lower, which means that the hot-rolling end temperature is set to 870°C or higher and 950°C or lower.
  • the thus-obtained hot-rolled steel sheet is wound at a temperature of 350°C or higher and 720°C or lower. Subsequently, the hot-rolled steel sheet is pickled and cold-rolled at a reduction ratio of 40% or higher and 90% or lower to obtain a cold-rolled steel sheet.
  • the hot-rolled steel sheet is produced through the typical steps of steel making, casting, and hot-rolling, but the hot-rolled steel sheet may be produced by thin slab casting without performing part or all of the hot-rolling steps.
  • the resultant cold-rolled steel sheet is annealed for 15 seconds or longer and 600 seconds or shorter in a first temperature range of 700°C or higher and 950°C or lower, specifically, in an austenite single-phase region or a dual-phase region of an austenite phase and a ferrite phase. If the annealing temperature is lower than 700°C or the annealing time is shorter than 15 seconds, a carbide in the steel sheet is sometimes not sufficiently dissolved, or the recrystallization of ferrite is not completed and thus desired ductility and stretch-flangeability are sometimes not achieved.
  • the annealing temperature is set in the range of 700°C or higher and 950°C or lower, preferably 760°C or higher and 920°C or lower.
  • the annealing time is set in the range of 15 seconds or longer and 600 seconds or shorter, preferably 30 seconds or longer and 400 seconds or shorter.
  • a second temperature range which is a temperature range from the first temperature range to 420°C
  • the annealed cold-rolled steel sheet is cooled to 550°C from the first temperature range at a cooling rate of 3 °C/s or higher, and is then cooled from 550 to 420°C within 600 seconds. Subsequently, the steel sheet is cooled at a cooling rate of 50 °C/s or lower in a third temperature range of 250°C or higher and 420°C or lower.
  • the cooling conditions in a second temperature range from the first temperature range to 420°C are essential to suppress the precipitation of phases other than intended ferrite and autotempered martensite phases.
  • the temperature range from the first temperature range to 550°C pearlite transformation easily occurs.
  • the cooling rate needs to be 3 °C/s or higher, and is preferably 5 °C/s or higher.
  • the upper limit of the cooling rate is not particularly specified, but special cooling equipment is required to achieve a cooling rate of 200 °C/s or higher.
  • the cooling rate is preferably 200 °C/s or lower.
  • the time required for cooling from 550°C to 420°C is 600 seconds or shorter, preferably 400 seconds or shorter.
  • the steel sheet is processed in the third temperature range.
  • autotempering treatment in which martensite transformation is caused while at the same time the transformed martensite is tempered is performed to obtain autotempered martensite in which the precipitation state of carbide grains is suitably controlled.
  • Typical martensite is obtained by performing annealing and then performing quenching with water cooling or the like.
  • the martensite is a hard phase, and contributes to an increase in the strength of a steel sheet but degrades workability.
  • a quenched steel sheet is normally heated again to perform tempering.
  • Fig. 1 schematically shows the steps described above. In such normal quenching and tempering treatments, after martensite transformation is completed by quenching, the temperature is increased to perform tempering. Consequently, a uniformly tempered microstructure is obtained.
  • autotempering treatment is a treatment in which a steel sheet is cooled in a certain cooling-rate range in the third temperature range as shown in Fig. 2 .
  • quenching and tempering through reheating are not performed, which is a method with high productivity.
  • the steel sheet including autotempered martensite obtained through this autotempering treatment has strength and workability equal to or higher than those of the steel sheet obtained by performing quenching and tempering through reheating shown in Fig. 1 .
  • martensite transformation and the tempering can be made to occur continuously or stepwise by performing continuous cooling (including stepwise cooling and holding) in the third temperature range. Consequently, a microstructure including martensites in different tempered states can be obtained.
  • the martensites in different tempered states have different characteristics in terms of strength and workability
  • desired characteristics as the entire steel sheet can be achieved by suitably controlling the amounts of martensites in different tempered states through autotempering treatment. Furthermore, since the autotempering treatment is performed without rapidly cooling a steel sheet to a low temperature range in which the martensite transformation is fully completed, the residual stress in the steel sheet is low and a steel sheet having a good plate shape is obtained, which is advantageous.
  • the third temperature range is 250°C or higher and 420°C or lower. If the temperature exceeds 420°C, bainite transformation is easily caused as described above. If the temperature is lower than 250°C, autotempering treatment requires a long time and thus proceeds insufficiently in a continuous annealing line or a continuous galvanizing and galvannealing line.
  • the cooling rate of a steel sheet needs to be 50 °C/s or lower in order to cause martensite transformation while at the same time tempering the transformed martensite and thus to obtain autotempered martensite. If the cooling rate exceeds 50 °C/s, the autotempering treatment insufficiently proceeds and the workability of martensite is sometimes not ensured. If the cooling rate is less than 0.1 °C/s, bainite transformation occurs or autotempering treatment excessively proceeds, whereby strength sometimes cannot be ensured. Thus, the cooling rate is preferably 0.1 °C/s or higher.
  • the steel sheet When a steel sheet is cooled at a cooling rate of 50 °C/s or lower in a third temperature range of 250°C or higher and 420°C or lower, the steel sheet is preferably cooled at a cooling rate of 1.0 °C/s or higher and 50 °C/s or lower in a temperature range of at least (Ms temperature - 50)°C or lower.
  • the area ratio of autotempered martensite in which the number of precipitated iron-based carbide grains each having a size of 0.1 ⁇ m or more and 0.5 ⁇ m or less is 5 ⁇ 10 2 or less per 1 mm 2 to the entire autotempered martensite can be set to 3% or more.
  • the cooling rate exceeds 50 °C/s, autotempering treatment insufficiently proceeds and desired autotempered martensite is not obtained. Consequently, the workability of martensite is sometimes not ensured. If the cooling rate is less than 1.0 °C/s, 3% or more of the area ratio of autotempered martensite in which the number of precipitated iron-based carbide grains each having a size of 0.1 ⁇ m or more and 0.5 ⁇ m or less is 5 ⁇ 10 2 or less per 1 mm 2 to the entire autotempered martensite cannot be achieved, and desired ductility and strength are not ensured. Thus, the cooling rate is set to 1.0 °C/s or higher.
  • Ms temperature can be obtained in a typical manner through the measurement of thermal expansion or electrical resistance during cooling. Alternatively, M obtained from an approximate expression (1) of Ms temperature described below may be used.
  • M °C 540 - 361 ⁇ C % / 1 - ⁇ % / 100 - 6 ⁇ Si % - 40 ⁇ Mn % + 30 ⁇ Al % - 20 ⁇ Cr % - 35 ⁇ V % - 10 ⁇ Mo % - 17 ⁇ Ni % - 10 ⁇ Cu %
  • [X%] is mass% of an alloy element X
  • [ ⁇ %] is the area ratio (%) of polygonal ferrite.
  • M represented by the above-described expression (1) is an empirical approximate expression of Ms temperature from which martensite transformation starts. It is believed that M is highly related to the precipitation behavior of iron-based carbide grains from martensite. Thus, M can be used as an indicator that indicates whether autotempered martensite in which the number of iron-based carbide grains each having a size of 5 nm or more and 0.5 ⁇ m or less is 5 ⁇ 10 4 or more per 1 mm 2 can be stably obtained. Even if M is less than 300°C, autotempered martensite is obtained. However, since the temperature is low, martensite transformation and autotempering treatment tend to slowly proceed. Compared with the case of M ⁇ 300°C, a steel sheet needs to be cooled slowly or held at a low temperature for a long time to obtain desired autotempered martensite, which may considerably lower manufacturing efficiency. Thus, M is preferably 300°C or higher.
  • the area ratio of polygonal ferrite is measured, for example, through the image processing and analysis of a SEM micrograph taken at 1000 to 3000 power. Polygonal ferrite is observed in the steel sheet that has been annealed and cooled under the above-described conditions. To ensure that M is 300°C or higher, after a cold-rolled steel sheet having a desired composition is produced, the area ratio of polygonal ferrite is measured and thus M is obtained from the expression (1) using the contents of alloy elements that can be calculated from the composition of the steel sheet. In the case where M is less than 300°C, the heat treatment conditions are suitably adjusted such that the area ratio of polygonal ferrite becomes lower, to obtain desired M. For example, the annealing temperature in the first temperature range is further increased and the average cooling rate from the first temperature range to 550°C is further increased. Alternatively, the contents of the components in the expression (1) may be adjusted.
  • the steel sheet of the present invention can be galvanized and galvannealed.
  • the galvanizing and galvannealing treatments are preferably performed in a continuous galvanizing and galvannealing line while the above-described annealing and cooling conditions are satisfied.
  • the galvanizing and galvannealing treatments are preferably performed in a temperature range of 420°C or higher and 550°C or lower.
  • the time required for cooling a steel sheet from 550°C to 420°C that is, the holding time in the temperature range of 420°C or higher and 550°C or lower needs to be 600 seconds or shorter, the time including galvanizing treatment time and/or galvannealing treatment time.
  • a method of galvanizing and galvannealing treatments is as follows. First, a steel sheet is immersed in a coating bath and the coating weight is adjusted using gas wiping or the like. In the case where the steel sheet is galvanized, the amount of dissolved Al in the coating bath is in the range of 0.12% or more and 0.22% or less. In the case where the steel sheet is galvannealed, the amount of dissolved Al is in the range of 0.08% or more and 0.18% or less.
  • the temperature of the coating bath is desirably 450°C or higher and 500°C or lower.
  • the temperature during alloying is preferably 450°C or higher and 550°C or lower. If the alloying temperature exceeds 550°C, an excessive amount of carbide grains are precipitated from untransformed austenite or the transformation into pearlite is caused, whereby desired strength and ductility are sometimes not achieved. Powdering is also degraded. If the alloying temperature is less than 450°C, the alloying does not proceed.
  • the coating weight is preferably in the range of 20 to 150 g/m 2 per surface. If the coating weight is less than 20 g/m 2 , corrosion resistance is degraded. Meanwhile, even if the coating weight exceeds 150 g/m 2 , the corrosion resistance is saturated, which merely increases the cost.
  • the degree of alloying is preferably in the range of 7 to 15% by mass on a Fe content basis in the coating layer. If the degree of alloying is less than 7% by mass, uneven alloying is caused and the surface appearance quality is degraded. Furthermore, a so-called ⁇ phase is formed in the coating layer and thus the slidability is degraded. If the degree of alloying exceeds 15% by mass, a large amount of hard brittle ⁇ phase is formed and the adhesion of the coating is degraded.
  • the holding temperature in the first temperature range, in the second temperature range, or the like is not necessarily constant. Even if the holding temperature is varied, the purport of the present invention is not impaired as long as the holding temperature is within a predetermined temperature range. The same is true for the cooling rate. Furthermore, a steel sheet may be subjected to annealing and autotempering treatments with any equipment as long as heat history is just satisfied. Moreover, it is also included in the scope of the present invention that, after autotempering treatment, temper rolling is performed on the steel sheet of the present invention for shape correction.
  • a slab to be formed into a steel sheet having the composition shown in Table 1 was heated to 1250°C and subjected to finish hot-rolling at 880°C.
  • the hot-rolled steel sheet was wound at 600°C, pickled, and cold-rolled at a reduction ratio of 65% to obtain a cold-rolled steel sheet having a thickness of 1.2 mm.
  • the resultant cold-rolled steel sheet was subjected to heat treatment under the conditions shown in Table 2. Quenching was not performed on any sample shown in Table 2.
  • the holding time in Table 2 was a time held at the holding temperature shown in Table 2.
  • the annealing time in a first temperature range of 700°C or higher and 950°C or lower was 600 seconds or shorter under any of the conditions shown in Table 2.
  • both surfaces were subjected to plating in a coating bath having a temperature of 463°C at a coating weight of 50 g/m 2 per surface.
  • the alloying treatment was performed such that Fe% (iron content) in the coating layer was adjusted to 9% by mass.
  • the resultant steel sheet was subjected to temper rolling at a reduction ratio (elongation ratio) of 0.3% regardless of the presence or absence of a coating.
  • the characteristics of the resultant steel sheets were evaluated by the following methods.
  • the area ratios of as-quenched martensite (untempered martensite) and retained austenite were obtained using the test pieces polished after heat treatment was performed at 200°C for 2 hours.
  • the test pieces polished after heat treatment was performed at 200°C for 2 hours were prepared in order to differentiate untempered martensite from retained austenite in the SEM observation. In the SEM observation, it is difficult to differentiate untempered martensite from retained austenite.
  • martensite is tempered, an iron-based carbide is formed in the martensite.
  • the iron-based carbide makes it possible to differentiate martensite from retained austenite.
  • the heat treatment at 200°C for 2 hours does not affect the phases other than martensite, that is, martensite can be tempered without changing the area ratio of each phase.
  • martensite can be differentiated from retained austenite due to the formed iron-based carbide.
  • the size and number of iron-based carbide grains included in autotempered martensite were measured through SEM observation.
  • the test pieces were the same as those used in the microstructure observation. Obviously, the test pieces polished without performing any treatment were observed. The test pieces were observed at a magnification of 10000x to 30000x in accordance with the precipitation state and size of the iron-based carbide grains.
  • the size of the iron-based carbide grains was evaluated using an average value of the major axis and minor axis of individual precipitates.
  • the number of iron-based carbide grains each having a size of 5 nm or more and 0.5 ⁇ m or less was counted and thus the number of iron-based carbide grains per 1 mm 2 of autotempered martensite was calculated. The observation was performed for 5 to 20 fields. The mean number was calculated from the total number of all the fields of each sample, and the mean number was employed as the number (per 1 mm 2 of autotempered martensite) of iron-based carbide grains of each sample.
  • the hardness HV of autotempered martensite was measured using an ultramicro-Vickers hardness meter at a load of 0.02 N. After the microstructure of autotempered martensite in which iron-based carbide grains were precipitated was confirmed by observing an indentation with a SEM, the average value of ten or more measurement values was employed as the hardness HV.
  • a tensile test was performed in accordance with JIS Z2241 using a JIS No. 5 test piece taken from the steel sheet in the rolling direction of the steel sheet.
  • Tensile strength (TS), yield strength (YS), and total elongation (T. El) were measured.
  • the product of the tensile strength and the total elongation (TS ⁇ T. El) was calculated to evaluate the balance between the strength and the elongation.
  • TS ⁇ T. E1 ⁇ 14500 MPa ⁇ %
  • Stretch-flangeability was evaluated in compliance with The Japan Iron and Steel Federation Standard JFST 1001.
  • the resulting steel sheet was cut into pieces each having a size of 100 mm ⁇ 100 mm.
  • a hole having a diameter of 10 mm was made in the piece by punching at a clearance of 12% of the thickness.
  • a cone punch with a 60° apex was forced into the hole while the piece was fixed with a die having an inner diameter of 75 mm at a blank-holding pressure of 88.2 kN. The diameter of the hole was measured when a crack was initiated.
  • Table 3 shows the evaluation results.
  • Table 3 Sample No. Steel type Area ratio (%) Mean hardness of autotempered martensite (HV) Number of iron-based carbide grains per 1 mm 2 *1 M (°C) YS (MPa) TS (Mpa) T. El (%) TS ⁇ T.
  • any steel sheet of the present invention has a tensile strength of 900 MPa or higher, a value of TS ⁇ T. El ⁇ 14500 (MPa ⁇ %), and a value of ⁇ ⁇ 15% that represents stretch-flangeability and thus has both high strength and good workability.
  • the steel sheets having an M of 300°C or higher are excellent in stretch-flangeability, particularly stretch-flangeability that is not degraded even if strength is increased.
  • the hardness of martensite is 700 ⁇ HV and the number of iron-based carbide grains included in martensite is less than 5 ⁇ 10 4 per 1 mm 2 or martensite does not include iron-based carbide grains. Therefore, a tensile strength of 900 MPa is satisfied, but a value of ⁇ is less than 15%, which provides poor workability. This is because, in sample Nos. 6 and 7, the cooling rate in the third temperature range is high, which does not satisfy 50 °C/s. In sample Nos. 3 and 8, the hardness of martensite is satisfactorily HV ⁇ 700, but the number of iron-based carbide grains included in martensite is less than 5 ⁇ 10 4 per 1 mm 2 .
  • a tensile strength of 900 MPa or higher is satisfied, but a value of ⁇ is less than 15%, which provides poor workability.
  • the cooling rate in the third temperature range is 55 °C/s, which does not satisfy 50 °C/s or lower.
  • TS ⁇ T. El is 14500 MPa ⁇ % or less.
  • the steel sheet of the present invention that includes autotempered martensite sufficiently subjected to autotempering treatment such that the hardness of martensite is HV ⁇ 700 and the number of iron-based carbide grains in martensite is 5 ⁇ 10 4 or more per 1 mm 2 has both high strength and good workability.
  • samples were manufactured in the same manner as the samples shown in Table 2, except that the cooling rate in a temperature range of 250°C or higher and (Ms temperature - 50)°C or lower of the third temperature range was changed as shown in Table. 4.
  • Table 4 sample Nos. 9, 11, 13, 14, and 26 are the same as those shown in Table 2 and listed in Table 4 to clarify the temperature range of 250°C or higher and (Ms temperature - 50)°C or lower. Note that M (°C) was used as the Ms temperature.
  • the characteristics of the thus-obtained steel sheets were evaluated in the same manner as in Example 1.
  • the amount of autotempered martensite in which the number of precipitated iron-based carbide grains each having a size of 0.1 ⁇ m or more and 0.5 ⁇ m or less is 5 ⁇ 10 2 or less per 1 mm 2 in the entire autotempered martensite was obtained as follows.
  • the test pieces polished without performing any treatment were observed at a magnification of 10000x to 30000x using a SEM.
  • the size of the iron-based carbide grains was evaluated using an average value of the major axis and minor axis of individual precipitates.
  • the area ratio of autotempered martensite in which the iron-based carbide grains have a size of 0.1 ⁇ m or more and 0.5 ⁇ m or less was measured. The observation was performed for 5 to 20 fields. Table 5 shows the results.

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KR101225321B1 (ko) 2013-01-22
WO2009096596A1 (ja) 2009-08-06
EP2246456A4 (de) 2014-04-23
EP2246456B1 (de) 2015-08-12
CN101932746A (zh) 2010-12-29
MX2010008227A (es) 2010-08-30
US20110030854A1 (en) 2011-02-10
CA2713181A1 (en) 2009-08-06
EP2246456B9 (de) 2016-12-21
CN101932746B (zh) 2012-05-09
JP5365217B2 (ja) 2013-12-11
JP2009203550A (ja) 2009-09-10
KR20100101697A (ko) 2010-09-17
CA2713181C (en) 2013-12-10

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