EP4685260A1 - High-strength steel sheet and method for producing same - Google Patents

High-strength steel sheet and method for producing same

Info

Publication number
EP4685260A1
EP4685260A1 EP24810653.6A EP24810653A EP4685260A1 EP 4685260 A1 EP4685260 A1 EP 4685260A1 EP 24810653 A EP24810653 A EP 24810653A EP 4685260 A1 EP4685260 A1 EP 4685260A1
Authority
EP
European Patent Office
Prior art keywords
less
steel sheet
steel
content
amount
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Pending
Application number
EP24810653.6A
Other languages
German (de)
French (fr)
Inventor
Ryohei Morimoto
Kazuki Endoh
Masaki Tada
Takeshi Nishiyama
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
JFE Steel Corp
Original Assignee
JFE Steel Corp
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by JFE Steel Corp filed Critical JFE Steel Corp
Publication of EP4685260A1 publication Critical patent/EP4685260A1/en
Pending legal-status Critical Current

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Classifications

    • 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
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • 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
    • C21D1/26—Methods of annealing
    • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/021—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving particular fabrication steps or treatments of ingots or slabs
    • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0273—Final recrystallisation annealing
    • 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
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • 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
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
    • 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
    • C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
    • 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
    • C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
    • 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
    • C22C38/10—Ferrous alloys, e.g. steel alloys containing cobalt
    • 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
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • 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
    • C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • 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
    • C23C2/04—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
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • 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
    • C21D2211/008—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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling

Definitions

  • the present invention relates to a steel sheet and a method for manufacturing the same.
  • the present invention relates to a high-strength steel sheet having excellent LME resistance, which is suitable as a member to be subjected to forming by cold pressing for use in an industrial field such as the automotive industry and the electrical machinery industry, and to a method for manufacturing the steel sheet.
  • the non-self-derived LME cracking is particularly noticeable when spot welding is performed with electrodes inclined at an angle.
  • the phrase "with electrodes inclined at an angle" herein refers to a state of welding electrodes in which the axis of the electrodes is not perpendicular to the surfaces of steel sheets.
  • Patent Literature 1 discloses a technique which involves controlling the frequency of corresponding grain boundaries in the surface of a steel sheet after a high-temperature tensile test and the thickness of a softened surface region, thereby achieving a high-strength steel sheet having a tensile strength of 980 MPa or more and a total elongation of 20% or more, and having excellent LME resistance.
  • Patent Literature 2 discloses a technique which involves introducing oxygen into the surface of a steel slab during continuous casting, thereby forming an iron oxide as a site for the formation of titanium nitride. This inhibits B in the steel from binding to dissolved nitrogen in a later annealing process and promotes the formation of (Fe,Mn) 2 B in a surface region of a steel sheet, thereby achieving a steel sheet having an LME resistance, high strength, and excellent ductility.
  • Patent Literature 3 proposes a technique which involves removing, prior to spot welding, a coating layer from a portion to be welded so as to prevent LME cracking.
  • Patent Literature 1 necessitates a reduction in the amount of Si in order to reduce the frequency of corresponding grain boundaries in the surface of a steel sheet. It therefore appears difficult to impart good formability to a steel sheet having a strength of 980 MPa or higher grade.
  • Patent Literature 2 which involves introducing oxygen into the surface of a steel slab during continuous casting to improve the LME resistance of the steel sheet finally obtained, it is expected that the introduction of oxygen into the surface of a slab will cause the formation of scale, resulting in a reduction in yield. Further, the presence of a large amount of hard particles (Fe,Mn) 2 B in a surface region from the casting stage will cause surface cracking during casting.
  • Patent Literature 3 requires a step of removing a coating layer in advance, which increases the production cost. Further, the removal of the coating layer will cause a reduction in the corrosion resistance of a weld.
  • the steel sheet include a hot-rolled steel sheet, a cold-rolled steel sheet, and a coated steel sheet such as GA or GI.
  • high strength refers to a TS of 980 MPa or more
  • good formability means that the relationship between tensile strength TS and elongation El, shown in the following formula 7, is satisfied.
  • the present inventors conducted intensive studies on the chemical compositions and microstructures of steel sheets. They have now found the following facts through precise control of slab heating conditions, temperatures from hot rolling to coiling, and annealing conditions, and through control of the states of elements existing in the steel.
  • the present invention is based on the above findings, and is summarized as follows.
  • the present invention it is possible to obtain a high-strength steel sheet having good formability and excellent LME resistance. Therefore, the present invention is very useful in the industrial fields of automobiles, electrical devices, etc., and is particularly useful for reducing the weight of automotive body frame parts.
  • the C content is an element necessary for increasing the strengths of tempered martensite, bainite, and fresh martensite. In order to fully achieve this effect, it is necessary to make the C content at least 0.030%. Therefore, the C content is made 0.030% or more.
  • the C content is preferably 0.050% or more, more preferably 0.070% or more, even more preferably 0.090% or more, and most preferably 0.100% or more.
  • the C content is made 0.500% or less.
  • the C content is preferably 0.400% or less, more preferably 0.300% or less, even more preferably 0.270% or less, and most preferably 0.250% or less.
  • Si more than 0.01% and not more than 2.50%
  • Si is an element that prevents excessive formation and growth of carbides in the steel, thereby increases the fraction of retained austenite and improving ductility. If the Si content is 0.01% or less, this effect is poor and the steel cannot have good formability. Therefore, the lower limit of the Si content is made 0.01%.
  • the Si content is made more than 0.01%.
  • the Si content is preferably 0.05% or more, more preferably 0.10% or more, even more preferably 0.50% or more, and most preferably 0.90% or more. However, if the Si content exceeds 2.50%, it causes a decrease in the melting point of zinc, which facilitates penetration of zinc into the steel sheet during welding, resulting in a reduction in the LME resistance of the steel sheet. Therefore, the Si content is made 2.50% or less.
  • the Si content is preferably 2.30% or less, more preferably 2.00% or less, even more preferably 1.80% or less, and most preferably 1.60% or less.
  • Mn not less than 0.10% and not more than 5.00%
  • Mn is an element that affects the area fractions of tempered martensite, bainite, and fresh martensite through enhancement of hardenability. If the Mn content is less than 0.10%, a soft phase such as ferrite will be formed excessively, and desired area fractions of tempered martensite, bainite, and fresh martensite cannot be obtained, resulting in an insufficient steel sheet strength. Therefore, the Mn content is made 0.10% or more.
  • the Mn content is preferably 0.50% or more, more preferably 0.80% or more, even more preferably 1.00% or more, and most preferably 2.00% or more.
  • the Mn content is made 5.00% or less.
  • the Mn content is preferably 4.50% or less, more preferably 4.00% or less, even more preferably 3.70% or less, and most preferably 3.50% or less.
  • the P content needs to be 0.100% or less.
  • the P content is preferably made 0.080% or less, more preferably 0.070% or less, even more preferably 0.050% or less, and most preferably 0.040% or less.
  • the P content is preferably made 0.001% or more, more preferably 0.003% or more, and even more preferably 0.005% or more.
  • the S content needs to be 0.0200% or less.
  • the S content is made 0.0200% or less.
  • the S content is preferably made 0.0180% or less, more preferably 0.0150% or less, even more preferably 0.0100% or less, and most preferably 0.0050% or less. While the lower limit of the S content is not particularly limited, in view of the restrictions of production technology, the S content is preferably 0.0001% or more, more preferably 0.0005% or more, and even more preferably 0.0010% or more.
  • Al is an element which acts as a deoxidizer and is effective in reducing inclusions in the steel, and is preferably added in a deoxidization process.
  • Al increases the transformation temperature for austenitizing, resulting in inclusion of ferrite in the microstructure. Therefore, the use of Al in an amount of more than 0.100% makes it difficult to achieve a desired TS.
  • the Al content is made 0.100% or less.
  • the Al content is preferably made 0.080% or less, more preferably 0.070% or less, even more preferably 0.060% or less, and most preferably 0.050% or less. While the lower limit of the Al content is not particularly limited, the Al content is preferably made 0.001% or more, more preferably 0.010% or more, and even more preferably 0.020% or more.
  • N may adversely affect the LME resistance through the formation of coarse nitrides. If the N content exceeds 0.0100%, a large amount of coarse nitrides will be formed, leading to a significant deterioration of LME resistance.
  • the N content is preferably as small as possible. Thus, the N content is made 0.0100% or less.
  • the N content is preferably 0.0090% or less, more preferably 0.0080% or less, even more preferably 0.0070% or less, and most preferably 0.0060% or less. While the lower limit of the N content is not particularly limited, in view of the restrictions of production technology, the N content is preferably made 0.0001% or more, more preferably 0.0010% or more, and even more preferably 0.0020% or more.
  • Ti contributes to precipitation strengthening and, in addition, reduces the prior austenite grain size and thereby reduces the grain sizes of tempered martensite and bainite.
  • Ti is effective in increasing the strength of the steel.
  • the Ti content is made 0.010% or more.
  • the Ti content is preferably 0.012% or more, more preferably 0.015% or more, even more preferably 0.020% or more, and most preferably 0.025% or more.
  • the Ti content is made 0.200% or less.
  • the Ti content is preferably 0.180% or less, more preferably 0.150% or less, even more preferably 0.100% or less, and most preferably 0.050% or less.
  • Nb not less than 0.005% and not more than 0.500%
  • Nb is an element that improves the LME resistance.
  • the presence of dissolved Nb in the steel during welding has a great effect on the improvement of the LME resistance.
  • the Nb content needs to be 0.005% or more.
  • the Nb content is preferably 0.007% or more, more preferably 0.008% or more, even more preferably 0.010% or more, and most preferably 0.012% or more.
  • the Nb content is made 0.500% or less.
  • the Nb content is preferably made 0.400% or less, more preferably 0.350% or less, even more preferably 0.200% or less, and most preferably 0.100% or less.
  • a high-strength steel sheet according to an embodiment of the present invention has a chemical composition containing the above-described components, with the balance including Fe and incidental impurities.
  • a high-strength steel sheet according to an embodiment of the present invention has a chemical composition containing the above-described components, with the balance consisting of Fe and incidental impurities.
  • the incidental impurities may include Zn, Pb, and As. Such impurities are allowed to be present in the steel in a total amount of up to 0.100%.
  • the chemical composition may contain, in mass %, one or two or more selected from the following: V: 0.500% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Mo: 1.00% or less, Co: 1.00% or less, Ni: 1.00% or less, Cu: 1.00% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.100% or less, Te: 0.100% or less, Hf: 0.10% or less, and Bi: 0.200% or less.
  • V 0.500% or less
  • V contributes to precipitation strengthening and, in addition, reduces the prior austenite grain size and thereby reduces the grain sizes of tempered martensite and bainite.
  • the steel can therefore contain V as necessary.
  • the lower limit of the V content is not particularly limited, in order to achieve the above effect, the V content is preferably made 0.001% or more, more preferably 0.005% or more, and even more preferably 0.010% or more.
  • the V content is made 0.500% or less.
  • the V content is preferably 0.400% or less, more preferably 0.300% or less, even more preferably 0.200% or less, and most preferably 0.100% or less.
  • Ta 0.10% or less
  • Ta like Ti, contributes to increasing the strength of the steel through the formation of an alloy carbide and an alloy carbonitride.
  • Ta partly dissolves in Nb carbide or Nb carbonitride to form a complex precipitate such as (Nb,Ta) (C,N), thereby significantly suppressing coarsening of precipitates and stabilizing the contribution of precipitation strengthening to the increase of the strength.
  • the steel can therefore contain Ta as necessary. While the lower limit of the Ta content is not particularly limited, in order to achieve the above effect, the Ta content is preferably made 0.01% or more, more preferably 0.02% or more, and even more preferably 0.03% or more. However, a too-high Ta content will lead to saturation of the precipitate stabilizing effect and increased alloying cost. Therefore, when the steel contains Ta, the Ta content is made 0.10% or less.
  • the Ta content is preferably made 0.08% or less, more preferably 0.07% or less, even more preferably 0.06% or less, and most preferably 0.05% or less.
  • the steel can contain W as necessary to improve the hardenability of the steel and to further increase the strength of the steel through a reduction in the grain sizes of tempered martensite and bainite.
  • the W content is preferably made 0.01% or more, more preferably 0.02% or more, and even more preferably 0.03% or more.
  • a W content of more than 0.10% may increase the amount of coarse precipitates, such as WN and WS, which remain undissolved during heating of the slab in a hot rolling process, resulting in a reduction in ductility. Therefore, when the steel contains W, the W content is made 0.10% or less.
  • the W content is preferably made 0.08% or less, more preferably 0.07% or less, even more preferably 0.06% or less, and most preferably 0.05% or less.
  • B is an element which, through its segregation at austenite grain boundaries, can improve hardenability.
  • B can form a microstructure composed mainly of tempered martensite and bainite, thereby increasing the strength of the steel sheet, and can contribute to improving the LME resistance. Therefore, the steel can contain B as necessary.
  • the lower limit of the B content is not particularly limited, in order to achieve the above effects, the B content is preferably made 0.0003% or more, more preferably 0.0005% or more, and even more preferably 0.0007% or more. However, if the B content exceeds 0.0100%, coarse precipitates will be formed, resulting in a reduction in ductility. Therefore, when the steel contains B, the B content is made 0.0090% or less.
  • the B content is preferably made 0.0080% or less, more preferably 0.0070% or less, even more preferably 0.0050% or less, and most preferably 0.0030% or less.
  • the steel can contain Cr as necessary. While the lower limit of the Cr content is not particularly limited, in order to achieve the above effect, the Cr content is preferably made 0.01% or more, more preferably 0.05% or more, and even more preferably 0.07% or more. However, if the steel contains Cr in an amount in excess of 1.00%, the area fraction of fresh martensite will be too high, resulting in a reduction in the dimensional accuracy and the ductility of the steel sheet during forming. Therefore, when the steel contains Cr, the Cr content is made 1.00% or less. The Cr content is preferably made 0.80% or less, more preferably 0.60% or less, even more preferably 0.50% or less, and most preferably 0.30% or less.
  • the steel can contain Mo as necessary. While the lower limit of the Mo content is not particularly limited, in order to achieve the above effect, the Mo content is preferably made 0.01% or more, more preferably 0.05% or more, and even more preferably 0.07% or more. However, if the steel contains Mo in an amount in excess of 1.00%, the area fraction of fresh martensite will be too high, resulting in a reduction in the dimensional accuracy and the ductility of the steel sheet during forming. Therefore, when the steel contains Mo, the Mo content is made 1.00% or less. The Mo content is preferably made 0.80% or less, more preferably 0.50% or less, even more preferably 0.30% or less, and most preferably 0.20% or less.
  • the steel can contain Co as necessary. While the lower limit of the Co content is not particularly limited, in order to achieve the above effect, the Co content is preferably made 0.01% or more, more preferably 0.05% or more, and even more preferably 0.07% or more. However, if the steel contains Co in an amount in excess of 1.00%, the area fraction of fresh martensite will be too high, resulting in a reduction in the dimensional accuracy and the ductility of the steel sheet during forming. Therefore, when the steel contains Co, the Co content is made 1.00% or less. The Co content is preferably made 0.80% or less, more preferably 0.60% or less, even more preferably 0.30% or less, and most preferably 0.20% or less.
  • Ni increases the strength of the steel through solid solution strengthening. Therefore, the steel can contain Ni as necessary. While the lower limit of the Ni content is not particularly limited, in order to achieve the above effect, the Ni content is preferably made 0.01% or more, more preferably 0.05% or more, and even more preferably 0.07% or more. However, if the steel contains Ni in an amount in excess of 1.00%, the area fraction of fresh martensite will be too high, resulting in a reduction in the dimensional accuracy and the ductility of the steel sheet during forming. Therefore, when the steel contains Ni, the Ni content is made 1.00% or less. The Ni content is preferably made 0.80% or less, more preferably 0.60% or less, even more preferably 0.30% or less, and most preferably 0.20% or less.
  • the steel can contain Cu as necessary. While the lower limit of the Cu content is not particularly limited, in order to achieve the above effect, the Cu content is preferably made 0.01% or more, more preferably 0.05% or more, and even more preferably 0.07% or more. However, if the steel contains Cu in an amount in excess of 1.00%, the area fractions of tempered martensite, bainite, and fresh martensite will be too high, resulting in a reduction in the dimensional accuracy and the ductility of the steel sheet during forming. Therefore, when the steel contains Cu, the Cu content is made 1.00% or less. The Cu content is preferably made 0.80% or less, more preferably 0.60% or less, even more preferably 0.30% or less, and most preferably 0.20% or less.
  • Sn and Sb suppress decarburization in a surface region having a thickness of about a few tens of ⁇ m, caused by nitridation or oxidation of the surface of the steel sheet, thereby preventing a reduction in the area fraction of tempered martensite in the steel sheet surface.
  • Sn and Sb have the effect of ensuring the strength and the material stability. Therefore, the steel can contain these elements as necessary. While the lower limit of the content of each element is not particularly limited, in order to achieve the above effects, the content of each element is preferably made 0.001% or more, more preferably 0.003% or more, and even more preferably 0.005% or more.
  • each element exceeds 0.200%, the ductility of the steel sheet may be reduced due to embrittlement of the steel sheet. Therefore, when the steel contains Sn and Sb, the content of each element is made 0.200% or less.
  • the content of each element is preferably 0.100% or less, more preferably 0.070% or less, even more preferably 0.050% or less, and most preferably 0.030% or less.
  • the content of each element is preferably made 0.0050% or less, more preferably 0.0040% or less, even more preferably 0.0035% or less, and most preferably 0.0030% or less.
  • each element makes the shape of a nitride or sulfide spherical and improves the limiting deformability of the steel sheet
  • the content of each of Ca, Mg and REM is preferably made 0.0001% or more, more preferably 0.0005% or more, even more preferably 0.0007% or more, and most preferably 0.0010% or more.
  • each in an amount of 0.100% or less causes no increase in the amount of coarse precipitates and inclusions and has no effect on the precipitation of Nb, and thus does not deteriorate the LME resistance. Therefore, when the steel contains Zr and Te, the content of each element is made 0.100% or less.
  • the content of each element is preferably made 0.080% or less, more preferably 0.070% or less, even more preferably 0.060% or less, and most preferably 0.050% or less.
  • each element makes the shape of a nitride or sulfide spherical and improves the limiting deformability of the steel sheet
  • the content of each of Zr and Te is preferably made 0.001% or more, more preferably 0.010% or more, and even more preferably 0.020% or more.
  • Hf in the steel in an amount of 0.10% or less causes no increase in the amount of coarse precipitates and inclusions and has no effect on the precipitation of Nb, and thus does not deteriorate the LME resistance. Therefore, when the steel contains Hf, the Hf content is made 0.10% or less.
  • the Hf content is preferably made 0.080% or less, more preferably 0.070% or less, even more preferably 0.060% or less, and most preferably 0.050% or less.
  • the Hf content is preferably made 0.003% or more, more preferably 0.010% or more, even more preferably 0.020% or more, and still more preferably 0.030% or more when the steel contains Hf.
  • the inclusion of Bi in the steel in an amount of 0.200% or less causes no increase in the amount of coarse precipitates and inclusions and has no effect on the precipitation of Nb, and thus does not deteriorate the LME resistance. Therefore, when the steel contains Bi, the Bi content is made 0.200% or less.
  • the Bi content is preferably made 0.100% or less, more preferably 0.050% or less, even more preferably 0.030% or less, and most preferably 0.020% or less. While the lower limit of the Bi content is not particularly specified, in view of the fact that Bi reduces segregation, the Bi content is preferably made 0.001% or more, more preferably 0.005% or more, and even more preferably 0.010% or more.
  • Tempered martensite and bainite contribute to the strength of the steel sheet.
  • the total area fraction of bainite and tempered martensite needs to be 40% or more.
  • the total area fraction is preferably 42% or more, more preferably 45% or more, even more preferably 47% or more, and most preferably 50% or more.
  • the steel microstructure is occupied by hard phases including tempered martensite, bainite, and fresh martensite.
  • the total area fraction of bainite and tempered martensite needs to be 85% or less.
  • the total area fraction is preferably 83% or less, more preferably 82% or less, even more preferably 81% or less, and most preferably 80% or less.
  • Fresh martensite is a very hard phase, and therefore increases the strength of the steel.
  • Fresh martensite is not necessarily required if the strength of the steel sheet is secured; however, the inclusion of fresh martensite in the steel sheet microstructure can further increase the strength of the steel sheet, achieving further strengthening. Therefore, the area fraction of fresh martensite needs to be 0% or more.
  • the area fraction is preferably 2% or more, more preferably 3% or more, even more preferably 4% or more, and most preferably 5% or more.
  • fresh martensite reduces the ductility of the steel, making it difficult to achieve good formability. Therefore, the area fraction of fresh martensite needs to be 25% or less.
  • the area fraction is preferably 23% or less, more preferably 22% or less, even more preferably 21% or less, and most preferably 20% or less.
  • the area fraction of retained austenite needs to be 5% or more.
  • the area fraction is preferably 7% or more, more preferably 8% or more, even more preferably 9% or more, and most preferably 10% or more.
  • the area fraction of retained austenite exceeds 20%, elements concentrated in the retained austenite, especially Si, will diffuse during welding, causing a decrease in the melting point of zinc. This may facilitate the penetration of zinc into the steel sheet during welding, resulting in a deterioration of the LME resistance of the steel sheet.
  • the area fraction of retained austenite needs to be 20% or less.
  • the area fraction is preferably 18% or less, more preferably 17% or less, even more preferably 16% or less, and most preferably 15% or less.
  • the inclusion of a remaining microstructure, such as ferrite or pearlite, in the steel does not impair the effects of the present invention.
  • a remaining microstructure it may include at least one of ferrite and pearlite, i.e., the area fraction of at least one of ferrite and pearlite may be 0% or more.
  • the area fraction of the remaining microstructure is preferably 20% or less, more preferably 18% or less, even more preferably 15% or less, and most preferably 13% or less.
  • Nb sol represents the amount (mass %) of dissolved Nb
  • Nb pre represents the amount (mass %) of Nb in Nb precipitates having a particle size of less than 20 nm.
  • Nb sol the amount of Nb in Nb precipitates having a particle size of less than 20 nm
  • Nb pre the amount of Nb in Nb precipitates having a particle size of less than 20 nm
  • Nb pre the amount of Nb in Nb precipitates having a particle size of less than 20 nm
  • Nb, unit: mass %) the total amount of Nb contained in the steel sheet
  • Nb in the steel sheet is brought to a dissolved state, or to a state which can be converted into a dissolved state during welding, thereby preventing non-self-derived LME and improving the LME resistance.
  • the present inventors through experiments conducted using various steel sheets with varying states of Nb present therein, have found that the LME resistance is improved when Formula 1 is satisfied.
  • Nb sol /Nb) + (Nb pre /Nb) is preferably 0.41 or more, more preferably 0.42 or more. While the upper limit is not particularly limited, "(Nb sol /Nb) + (Nb pre /Nb)” is preferably 0.95 or less, more preferably 0.90 or less.
  • the lower limit of the particle size of the Nb precipitates is not particularly limited; for example, the Nb precipitates may have a particle size of 0.1 nm or more.
  • Amount of Diffusible Hydrogen in Steel 0.50 mass ppm or less
  • the amount of diffusible hydrogen in the steel is made 0.50 mass ppm or less.
  • the amount of diffusible hydrogen is preferably made 0.30 mass ppm or less, more preferably 0.25 mass ppm or less, even more preferably 0.20 mass ppm or less, and most preferably 0.15 mass ppm or less. While the lower limit of the amount of diffusible hydrogen in the steel is not particularly specified, in view of the restrictions of production technology, the amount of diffusible hydrogen is preferably made 0.01 mass ppm or more, more preferably 0.02 mass ppm or more, even more preferably set to 0.03 mass ppm or more, and most preferably 0.05 mass ppm or more.
  • the steel material heating temperature needs to be equal to or higher than T sol °C. Therefore, the steel material heating temperature is made equal to or higher than T sol °C represented by the following Formula 2.
  • the steel material heating temperature is preferably equal to or higher than "T sol ⁇ 1.1"°C, more preferably equal to or higher than "T sol ⁇ 1.2"°C, even more preferably equal to or higher than "T sol ⁇ 1.3"°C, and most preferably equal to or higher than "T sol ⁇ 1.5"°C.
  • the heating time at the temperature is made 1.0 hours or more.
  • the finish rolling start temperature is preferably made "T sol °C - 100°C” or higher, more preferably “T sol °C - 70°C” or higher, even more preferably “T sol °C - 50°C” or higher, and most preferably "T sol °C - 30°C” or higher.
  • a finish rolling start temperature exceeding 1200°C may increase a scale loss during pre-heating of the slab, causing breakage of the sheet during hot rolling. Therefore, the finish rolling start temperature in hot rolling is preferably made 1200°C or lower, more preferably 1170°C or lower, even more preferably 1150°C or lower, and most preferably 1120°C or lower.
  • the steel material after the heating is hot-rolled into a hot-rolled steel sheet. While the upper and lower limits of the delivery temperature in the finish rolling are not particularly specified, a finish rolling delivery temperature of less than 800°C will cause precipitation of Nb which has been dissolved during heating of the steel material, resulting in an increase in the amount of Nb precipitated. Further, such a delivery temperature increases a rolling load, and thus a rolling burden, which may impede a cold rolling process. Therefore, the finish rolling delivery temperature in hot rolling is preferably made 800°C or higher, more preferably 820°C or higher, even more preferably 840°C or higher, and most preferably 850°C or higher.
  • the finish rolling delivery temperature is preferably made 1000°C or lower, more preferably 980°C or lower, even more preferably 970°C or lower, and most preferably 950°C or lower.
  • Effective time t HR is herein defined as the time required from the start of finish rolling to the end of finish rolling. Since Nb contained in the steel sheet begins to precipitate and grow during finish rolling, the effective time t HR is one of the parameters for controlling the form of Nb to improve the LME resistance. While the upper and lower limits of the effective time t HR are not particularly specified, the effective time t HR is preferably made 3 seconds or more, more preferably 4 seconds or more, even more preferably 5 seconds or more, and most preferably 7 seconds or more. The effective time t HR is preferably made 15 seconds or less, more preferably 12 seconds or less, even more preferably 11 seconds or less, and most preferably 10 seconds or less.
  • Residence time t CT is herein defined as the time from the finish rolling delivery temperature to 650°C. Since the precipitation of Nb and the growth of precipitated Nb are likely to proceed during the period from the completion of finish rolling to the sheet temperature reaching 650°C, the residence time t CT is one of the parameters for controlling the form of Nb to improve the LME resistance. While the upper and lower limits of the residence time t CT are not particularly specified, the residence time t CT is preferably made 5 seconds or more, more preferably 7 seconds or more, even more preferably 8 seconds or more, and most preferably 9 seconds or more. The residence time t CT is made 20 seconds or less, more preferably 18 seconds or less, even more preferably 17 seconds or less, and most preferably 15 seconds or less.
  • the coiling temperature after hot rolling is made 650°C or lower.
  • the coiling temperature is preferably 630°C or lower, more preferably 620°C or lower, even more preferably 610°C or lower, and most preferably 600°C or lower.
  • the lower limit of the coiling temperature is not particularly limited, if the coiling temperature is lower than 300°C, the strength of the hot-rolled sheet increases, which may cause an increase in the rolling burden in cold rolling and a defective sheet shape, resulting in a reduction in productivity. Therefore, the lower limit of the coiling temperature is preferably made 300°C or higher, more preferably 320°C or higher, even more preferably 350°C or higher, and most preferably 400°C or higher.
  • the resulting hot-rolled steel sheet may be subjected to intermediate heat treatment at a temperature below 650°C, as necessary, to prevent an increase in load in later cold rolling.
  • the heat treatment is preferably performed at a temperature of 150°C or higher.
  • the heat treatment may be carried out, for example, in a box annealing furnace at a soaking temperature of 500°C for a soaking time of 4 hours.
  • the resulting hot-rolled steel sheet may be subjected to a treatment, such as pickling, as necessary. Pickling of the hot-rolled coil may be performed by a common method.
  • the hot-rolled coil may also be subjected to skin pass rolling to correct its shape and enhance its pickling properties.
  • the steel sheet may be subjected directly to the below-described annealing step (heat treatment), or may be subjected to cold rolling prior to the heat treatment.
  • the cold rolling reduction is preferably made 25% or more, more preferably 30% or more, even more preferably 32% or more, and most preferably 35% or more.
  • an excessive reduction causes a too-high rolling load, leading to an increase in the burden on a cold rolling mill. Therefore, the cold rolling reduction is preferably made 75% or less, more preferably 70% or less, even more preferably 67% or less, and most preferably 65% or less.
  • the heating time t is one of the parameters for controlling the form of Nb to improve the LME resistance. While the upper and lower limits of the heating time t are not particularly specified, the heating time t is preferably made 300 seconds or more, more preferably 400 seconds or more, even more preferably 450 seconds or more, and most preferably 490 seconds or more. The heating time t is preferably made 700 seconds or less, more preferably 650 seconds or less, even more preferably 600 seconds or less, and most preferably 590 seconds or less.
  • Holding time t AT at the soaking temperature T AT °C is one of the control parameters for promoting austenitizing of the steel sheet during soaking and, since it is related to the precipitation of Nb and the growth of Nb precipitates, is also one of the parameters for controlling the LME resistance. While the upper and lower limits of the holding time t AT are not particularly specified, the holding time t AT is preferably made 15 seconds or more in order to sufficiently promote austenitizing. The holding time t AT is more preferably made 30 seconds or more, even more preferably 50 seconds or more, and most preferably 100 seconds or more.
  • Q HR 0.5 T FET + T FDT ⁇ log 10 t HR
  • Q CT 0.5 T CT + 650 ⁇ log 10 t CT
  • Q AT 0.5 650 + T AT ⁇ log 10 t + T AT ⁇ log 10 t AT Q HR + Q CT + Q AT ⁇ 6000
  • Q HR , Q CT , and Q AT are each defined by the temperatures at the start and end of each of the effective time t HR , the residence time t CT , the heating time t, and the holding time t AT , during which heat of 650°C to T sol is applied.
  • the sum of Q HR , Q CT , and Q AT is made 6000 or less, preferably 5990 or less, more preferably 5980 or less, even more preferably 5970 or less, and most preferably 5950 or less.
  • Q AT is preferably made 4700 or less, more preferably 4650 or less, even more preferably 4630 or less, and most preferably 4600 or less.
  • the steel sheet may be reheated.
  • C is concentrated in untransformed austenite. This enhances the stability of austenite and increases the area fraction of retained austenite contained in the steel sheet after cooling, making it possible to further enhance the ductility.
  • the reheating temperature is preferably made 200°C or higher, more preferably 210°C or higher, even more preferably 230°C or higher, and most preferably 250°C or higher.
  • the reheating temperature is preferably made 450°C or lower, more preferably 430°C or lower, even more preferably 410°C or lower, and most preferably 400°C or lower.
  • the coating weight (per one surface) is preferably 20 g/m 2 or more from the viewpoint of corrosion resistance and control of the coating weight.
  • the coating weight is more preferably 25 g/m 2 or more, even more preferably 30 g/m 2 or more, and most preferably 32 g/m 2 or more.
  • the coating weight is preferably 120 g/m 2 or less, more preferably 100 g/m 2 or less, even more preferably 70 g/m 2 or less, and most preferably 65 g/m 2 or less.
  • the hot-dip galvanizing is preferably performed using a galvanizing bath containing 0.08% to 0.30% of Al.
  • the amount of Al in the hot-dip galvanizing is preferably made 0.08% or more, more preferably 0.09% or more, even more preferably 0.10% or more, and most preferably 0.12% or more.
  • the amount of Al in the hot-dip galvanizing is preferably made 0.30% or less, more preferably 0.25% or less, even more preferably 0.22% or less, and most preferably 0.20% or less.
  • the treatment is performed in a temperature range of 450°C to 600°C. If the alloying treatment is performed at a temperature exceeding 600°C, untransformed austenite will be transformed into pearlite, and the area fraction of retained austenite will be less than 5%, which may result in a reduction in ductility. Therefore, when an alloying treatment of a galvanized coating is performed, the alloying treatment is preferably performed in a temperature range of not less than 450°C, more preferably not less than 460°C, even more preferably not less than 465°C, and most preferable not less than 470°C.
  • the alloying treatment is preferably performed in a temperature range of not more than 600°C, more preferably not more than 570°C, even more preferably not more than 550°C, and most preferably not more than 530°C.
  • the alloyed coating layer of the hot-dip galvanized steel sheet preferably has an Fe concentration of 8% to 17%.
  • the Fe concentration of the alloyed coating layer of the hot-dip galvanized steel sheet is preferably 8% or more, more preferably 9% or more, and even more preferably 10% or more.
  • the Fe concentration of the alloyed coating layer of the hot-dip galvanized steel sheet is preferably 17% or less, more preferably 16% or less, and even more preferably 15% or less.
  • the alloyed coating layer is thus formed by performing the alloying treatment on the hot-dip galvanized steel sheet.
  • the steel sheet after annealing was held at a holding temperature of 200°C to 450°C for 10 seconds or more, and then cooled to room temperature. In either case, the cooled steel sheet was cold-rolled at a rolling reduction of 50% to obtain a high-strength cold-rolled steel sheet (CR).
  • CR high-strength cold-rolled steel sheet
  • HR hot-rolled steel sheets
  • G hot-dip galvanized steel sheets
  • GA galvannealed steel sheets
  • a galvanizing bath containing 0.19 mass % Al was used for the hot-dip galvanized steel sheets (GI), and a galvanizing bath containing 0.14 mass % Al was used for the galvannealed steel sheets (GA).
  • the temperature of each bath was 465°C.
  • the coating weight was 45 g/m 2 per one surface (double-sided coating).
  • the Fe concentration in the coating layer was adjusted to be within the range of 9 mass % to 12 mass %.
  • TM tempered martensite
  • B bainite
  • FM fresh martensite
  • yR retained austenite
  • the area fractions of fresh martensite, tempered martensite, and bainite were determined by the following method.
  • the cross-section was observed by a scanning electron microscope (SEM) at 2000-fold magnification in 10 fields of view at a 1/4 thickness position (a position at a distance of 1/4 of the sheet thickness from the steel sheet surface).
  • SEM scanning electron microscope
  • the area fractions of constituent microstructures (the total fraction of tempered martensite and bainite, and the fraction of fresh martensite) were calculated.
  • fresh martensite was determined by a light gray microstructural region
  • tempered martensite and bainite were determined by a dark gray region in which carbides were precipitated.
  • the area fraction of retained austenite was determined by the following method. Each steel sheet was polished from the 1/4 thickness position to a plane by a thickness of 0.1 mm, and then further chemically polished to a plane by a thickness of 0.1 mm. The plane (surface) was subjected to X-ray diffraction analysis using CoK ⁇ rays to measure the integrated intensity ratios of the diffraction peaks of the ⁇ 200 ⁇ , ⁇ 220 ⁇ and ⁇ 311 ⁇ planes of fcc iron, and the ⁇ 200 ⁇ , ⁇ 211 ⁇ and ⁇ 220 ⁇ planes of bcc iron. The nine integrated intensity ratios obtained were averaged to determine the area fraction of retained austenite.
  • the amount of hydrogen in steel was determined by the following method. Test specimens of about 5 ⁇ 30 mm were cut out from a hot-rolled steel sheet, a cold-rolled steel sheet, and a galvanized steel sheet. For the test specimen of galvanized steel sheet, the coating on the surface of the test specimen was removed in advance using a router (precision grinder). Each test specimen was placed in a quartz tube, and the internal atmosphere of the tube was replaced with Ar, and then heated at 200°C/hr. Hydrogen generated until the temperature reached 400°C was subjected to gas chromatography to measure the amount of released hydrogen by a temperature rise analysis method. The cumulative value of the amount of hydrogen detected in the temperature range from room temperature (25°C) to less than 250°C was taken as the amount of diffusible hydrogen.
  • the total amount of Nb contained in the steel (Nb total ), the amount of dissolved Nb in the steel sheet (Nb sol ), and the amount of Nb in Nb precipitates having a particle size of less than 20 nm (Nb pre ) were measured in the following manner.
  • the total amount of Nb contained in the steel (Nb) was measured by wet chemical analysis.
  • the amount of Nb in Nb precipitates having a particle size of 20 nm or less was measured by the following method. Precipitates in the steel material were captured as residues, and the amount of Nb in all residues was determined. Thereafter, the amount of Nb present in residues having a particle size of 20 nm or more was determined.
  • the amount of Nb in Nb precipitates having a particle size of 20 nm or less was determined as the difference between the amount of Nb in all residues and the amount of Nb present in the residues having a particle size of 20 nm or more. Specific procedures are as follows. A plurality of test specimens of about 20 ⁇ 50 mm were cut out from a hot-rolled steel sheet, a cold-rolled steel sheet, or a galvanized steel sheet. For a test specimen of a galvanized steel sheet, the coating on the surface of the test specimen was removed in advance using a router (precision grinder). The surface of each test specimen was polished about 50 ⁇ m by preliminary electrolytic polishing to obtain a fresh surface.
  • the resulting test specimen was subjected to electrolysis using 10 vol % acetylacetone - 1 mass % tetramethylammonium chloride - methanol as an electrolytic solution for extracting precipitates.
  • the resulting electrolytic solution after electrolysis was passed through a filter having a pore size of 0.2 ⁇ m to capture residues.
  • the residues were then decomposed with an acid, and the Nb concentration was quantified in mass % unit using ICP emission spectrometry. The value obtained was used as the amount of Nb in all residues.
  • a remaining test specimen was subjected to electrolysis using 10 vol % acetylacetone - 1 mass % tetramethylammonium chloride - methanol.
  • the residues captured on the alumina filter had a particle size of 20 nm or more.
  • the residues were decomposed with an acid, and the Nb concentration was quantified in mass % unit using ICP emission spectrometry.
  • the value obtained was used as the amount of Nb present in the residues having a particle size of 20 nm or more.
  • the difference between the amount of Nb in all residues and the amount of Nb present in the residues having a particle size of 20 nm or more was taken as the amount of Nb in Nb precipitates having a particle size of 20 nm or less (Nb pre ).
  • the amount of dissolved Nb in the steel sheet (Nb sol ) was determined as the difference between the total amount of Nb contained in the steel (Nb) and the amount of Nb in all residues determined by the above method.
  • Nb precipitates are composed mainly of NbC
  • an NbN precipitate or other Nb precipitates may be contained in the steel sheet.
  • a tensile test was performed according to JIS Z 2241 (2011) using a JIS No. 5 test specimen which had been taken from a steel sheet such that the tensile direction was perpendicular to the rolling direction of the steel sheet, and the TS (tensile strength) and EL (total elongation) of the test specimen were measured.
  • steel sheets having a TS of 980 MPa or more were evaluated as acceptable, judging that the intended strength was achieved.
  • Steel sheets having a "TS 1.5 ⁇ El" of 390,000 or more were evaluated as acceptable, judging that the intended formability was achieved.
  • the LME resistance was evaluated in the following manner.
  • a sample having the dimensions of: 100 mm in a direction perpendicular to the rolling direction, and 30 mm in the rolling direction, was taken from a steel sheet.
  • the sample and a 980GA sample having the same size were stacked to prepare an evaluation sample.
  • Resistance spot welding was performed on the evaluation sample under the conditions of an electrode inclination angle of ⁇ and a welding pressure of 3.5 kN.
  • the electrode inclination angle in spot welding is herein defined as the angle ⁇ between a line passing through the major axis of a nugget and a line parallel to the surface of a steel sheet in a cross-section of a spot-welded member.
  • the welding current pattern was controlled so that the diameter of a nugget formed would fall within the range of 3.5 ⁇ t to 5.5 ⁇ t, t being the thickness (1.2 mm) of one steel sheet.
  • Dr6-type CuCr electrodes were used, and the clearance between the evaluation sample and an electrode was 2.0 mm.
  • the electrode inclination angles and holding times in the LME resistance evaluation tests are shown in Tables 3-1 and 3-2.
  • the high-strength steel sheets of the Inventive Examples all have a high TS and a good TS ⁇ El balance, and are excellent in LME resistance, whereas the steel sheets of the Comparative Examples are poor in at least one of TS, TS ⁇ El balance, and LME resistance.

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Abstract

The present invention is directed to providing a steel sheet having excellent LME resistance, high strength, and good formability by a method different from conventional techniques, and to providing a method for manufacturing the steel sheet.
A high-strength steel sheet containing particular components and having a particular microstructure, wherein the relationship between the amount of dissolved Nb (Nbsol), the amount of Nb in Nb precipitates having a particle size of less than 20 nm (Nbpre), and the total amount of Nb contained in the steel sheet (Nb) satisfies the following Formula 1, and wherein the amount of diffusible hydrogen in the steel is 0.50 mass ppm or less: Nb sol / Nb + Nb pre / Nb ≥ 0.40 where Nbsol represents the amount (mass %) of dissolved Nb, and Nbpre represents the amount (mass %) of Nb in Nb precipitates having a particle size of less than 20 nm.

Description

    Technical Field
  • The present invention relates to a steel sheet and a method for manufacturing the same. In particular, the present invention relates to a high-strength steel sheet having excellent LME resistance, which is suitable as a member to be subjected to forming by cold pressing for use in an industrial field such as the automotive industry and the electrical machinery industry, and to a method for manufacturing the steel sheet.
    • Background Art
  • In recent years, from the viewpoint of global environmental protection, there has been an increasing need for weight reduction of car bodies to improve fuel efficiency. Therefore, the use of high-strength steel sheets for automotive parts is desired. When cold-rolled steel sheets are used for automotive parts, a method is generally used in which two or more steel sheets, which have undergone a forming process, are joined by welding and finished into a desired shape.
  • In the production of automotive parts, spot welding is mainly used to join steel sheets together. It is known that when welding is performed on a sheet assembly including a high-strength steel sheet (high-strength galvanized steel sheet) having a welded galvanizing layer, cracking is likely to occur in a weld.
  • This is considered to be due to the fact that a galvanized coating layer, which generally has a low melting point, will liquefy during welding and penetrate into grain boundaries in the base steel sheet, causing a reduction in the grain boundary strength. Such a reduction in strength due to liquefied metal is generally called liquid metal embrittlement (LME). Cracking that occurs due to liquid metal embrittlement is called liquid metal embrittlement (LME) cracking.
  • It is also known that cracking occurs in a weld even in a high-strength steel sheet having no coating when a mating material to which the high-strength steel sheet is welded has a galvanized coating layer. Such LME cracking that occurs in a high-strength steel sheet having no coating layer is herein referred to as non-self-derived LME cracking.
  • The non-self-derived LME cracking is particularly noticeable when spot welding is performed with electrodes inclined at an angle. The phrase "with electrodes inclined at an angle" herein refers to a state of welding electrodes in which the axis of the electrodes is not perpendicular to the surfaces of steel sheets.
  • Various studies have been conducted to prevent LME cracking in a galvanized steel sheet and non-self-derived LME cracking in a non-coated steel sheet upon spot welding of the steel sheets.
  • For example, Patent Literature 1 discloses a technique which involves controlling the frequency of corresponding grain boundaries in the surface of a steel sheet after a high-temperature tensile test and the thickness of a softened surface region, thereby achieving a high-strength steel sheet having a tensile strength of 980 MPa or more and a total elongation of 20% or more, and having excellent LME resistance.
  • Patent Literature 2 discloses a technique which involves introducing oxygen into the surface of a steel slab during continuous casting, thereby forming an iron oxide as a site for the formation of titanium nitride. This inhibits B in the steel from binding to dissolved nitrogen in a later annealing process and promotes the formation of (Fe,Mn)2B in a surface region of a steel sheet, thereby achieving a steel sheet having an LME resistance, high strength, and excellent ductility.
  • Patent Literature 3 proposes a technique which involves removing, prior to spot welding, a coating layer from a portion to be welded so as to prevent LME cracking.
    • Citation List Patent Literature
    • Summary of Invention Technical Problem
  • However, the technique described in Patent Literature 1 necessitates a reduction in the amount of Si in order to reduce the frequency of corresponding grain boundaries in the surface of a steel sheet. It therefore appears difficult to impart good formability to a steel sheet having a strength of 980 MPa or higher grade.
  • Regarding the technique described in Patent Literature 2 which involves introducing oxygen into the surface of a steel slab during continuous casting to improve the LME resistance of the steel sheet finally obtained, it is expected that the introduction of oxygen into the surface of a slab will cause the formation of scale, resulting in a reduction in yield. Further, the presence of a large amount of hard particles (Fe,Mn)2B in a surface region from the casting stage will cause surface cracking during casting.
  • The technique described in Patent Literature 3 requires a step of removing a coating layer in advance, which increases the production cost. Further, the removal of the coating layer will cause a reduction in the corrosion resistance of a weld.
  • It is therefore an object of the present invention to provide a steel sheet having excellent LME resistance, high strength, and good formability by a method different from the conventional techniques, and to provide a method for manufacturing the steel sheet. Examples of the steel sheet include a hot-rolled steel sheet, a cold-rolled steel sheet, and a coated steel sheet such as GA or GI.
  • In the present invention, "high strength" refers to a TS of 980 MPa or more, and "good formability" means that the relationship between tensile strength TS and elongation El, shown in the following formula 7, is satisfied. TS 1.5 × El ≥ 390000
    • Solution to Problem
  • In order to solve the above problems, the present inventors conducted intensive studies on the chemical compositions and microstructures of steel sheets. They have now found the following facts through precise control of slab heating conditions, temperatures from hot rolling to coiling, and annealing conditions, and through control of the states of elements existing in the steel.
  • By dissolving Nb or by controlling the particle radius of precipitates within a particular range which will not cause cracking on a steel surface during casting, a dissolved element which is present in the steel during welding or a dissolved element produced by dissolution of the precipitates prevents penetration of zinc into the steel, thereby improving the LME resistance. On the other hand, when the radius of precipitates of a particular element is larger than a predetermined value, the precipitates will not be dissolved sufficiently during welding, resulting in a failure to achieve an improvement in LME resistance by the dissolved element. In addition, it is quite likely that cracking will newly occur starting from coarse precipitates not only during a manufacturing process but after press forming or during welding.
  • The present invention is based on the above findings, and is summarized as follows.
    1. [1] A high-strength steel sheet having a chemical composition containing, in mass %:
      • C: not less than 0.030% and not more than 0.500%,
      • Si: more than 0.01% and not more than 2.50%,
      • Mn: not less than 0.10% and not more than 5.00%,
      • P: 0.100% or less,
      • S: 0.0200% or less,
      • Al: 0.100% or less,
      • N: 0.0100% or less,
      • O: 0.0100% or less,
      • Ti: not less than 0.010% and not more than 0.200%, and
      • Nb: not less than 0.005% and not more than 0.500%,
      • with the balance being Fe and incidental impurities,
        • wherein in the microstructure of the steel sheet at a 1/4 thickness position,
        • the total area fraction of tempered martensite and bainite is not less than 40% and not more than 85%,
        • the area fraction of fresh martensite is not less than 0% and not more than 25%,
        • the area fraction of retained austenite is not less than 5% and not more than 20%, and
        • the balance includes at least one of ferrite and pearlite at an area fraction of not less than 0% and not more than 20%,
        • wherein the relationship between the amount of dissolved Nb (Nbsol), the amount of Nb in Nb precipitates having a particle size of less than 20 nm (Nbpre), and the total amount of Nb contained in the steel sheet (Nb) satisfies the following Formula 1, and
        • wherein the amount of diffusible hydrogen in the steel is 0.50 mass ppm or less: Nb sol / Nb + Nb pre / Nb ≥ 0.40 where Nbsol represents the amount (mass %) of dissolved Nb, and Nbpre represents the amount (mass %) of Nb in Nb precipitates having a particle size of less than 20 nm.
    2. [2] The high-strength steel sheet according to [1], wherein the chemical composition further contains, in mass %, one or two or more selected from the following:
      • V: 0.500% or less,
      • Ta: 0.10% or less,
      • W: 0.10% or less,
      • B: 0.0100% or less,
      • Cr: 1.00% or less,
      • Mo: 1.00% or less,
      • Co: 1.00% or less,
      • Ni: 1.00% or less,
      • Cu: 1.00% or less,
      • Sn: 0.200% or less,
      • Sb: 0.200% or less,
      • Ca: 0.0100% or less,
      • Mg: 0.0100% or less,
      • REM: 0.0100% or less,
      • Zr: 0.100% or less,
      • Te: 0.100% or less,
      • Hf: 0.10% or less, and
      • Bi: 0.200% or less.
    3. [3] The high-strength steel sheet according to [1] or [2], having a coating layer on the surface of the steel sheet.
    4. [4] The high-strength steel sheet according to [3], wherein the coating layer is an alloyed coating layer.
    5. [5] A method for manufacturing the high-strength steel sheet according to [1] or [2], including:
      • a slab heating step of heating a steel material having the chemical composition at a temperature equal to or higher than Tsol°C represented by the following Formula 2 for 1.0 hours or more;
      • a hot rolling step;
      • a coiling step at a coiling temperature TCT°C of 650°C or lower; and
      • an annealing step of heating a steel sheet to a soaking temperature TAT°C, holding the steel sheet at the soaking temperature TAT°C of not less than 750°C and not more than 950°C, and then cooling the steel sheet,
      • wherein QHR defined by Formula 3 in terms of a finish rolling start temperature TFET°C, a finish rolling delivery temperature TFDT°C, and an effective time tHR from the start of finish rolling to the end of finish rolling in the hot rolling step,
      • QCT defined by Formula 4 in terms of a residence time tCT from TFDT°C to 650°C and the coiling temperature TCT°C in the coiling step, and
      • QAT defined by Formula 5 in terms of the soaking temperature TAT°C, a heating time t from 650°C to the soaking temperature TAT°C, and a holding time tAT at the soaking temperature TAT°C in the annealing step, satisfy Formula 6: T sol = 7900 / 3.42 − log Nb % C % − 273 where [Nb%] and [C%] respectively represent the Nb content (mass %) and the C content (mass %) of the steel; Q HR = 0.5 T FET + T FDT × log 10 t HR ; Q CT = 0.5 T CT + 650 × log 10 t CT ; QAT = 0.5(650 + TAT) × log10(t) + TAT × log10(tAT); Q HR + Q CT + Q AT ≤ 6000 .
    6. [6] The method for manufacturing a high-strength steel sheet according to [5], wherein the value of QAT is 4700 or less.
    7. [7] The method for manufacturing a high-strength steel sheet according to [5] or [6], wherein the steel sheet after the annealing step is subjected to a coating treatment.
    8. [8] The method for manufacturing a high-strength steel sheet according to [7], wherein the coating treatment is an alloying coating treatment.
    Advantageous Effects of Invention
  • According to the present invention, it is possible to obtain a high-strength steel sheet having good formability and excellent LME resistance. Therefore, the present invention is very useful in the industrial fields of automobiles, electrical devices, etc., and is particularly useful for reducing the weight of automotive body frame parts.
    • Description of Embodiments
  • The present invention will now be described in detail. In the following description of component elements, "%" refers to "mass %" unless otherwise specified.
    1. (1) The following are the reasons for the above-described limitations on the chemical composition of the steel of the present invention.
    C: not less than 0.030% and not more than 0.500%
  • C is an element necessary for increasing the strengths of tempered martensite, bainite, and fresh martensite. In order to fully achieve this effect, it is necessary to make the C content at least 0.030%. Therefore, the C content is made 0.030% or more. The C content is preferably 0.050% or more, more preferably 0.070% or more, even more preferably 0.090% or more, and most preferably 0.100% or more. On the other hand, if the C content exceeds 0.500%, the weldability and the LME resistance of the steel, which are important in joining automotive parts, will deteriorate. Therefore, the C content is made 0.500% or less. The C content is preferably 0.400% or less, more preferably 0.300% or less, even more preferably 0.270% or less, and most preferably 0.250% or less.
    • Si: more than 0.01% and not more than 2.50%
  • Si is an element that prevents excessive formation and growth of carbides in the steel, thereby increases the fraction of retained austenite and improving ductility. If the Si content is 0.01% or less, this effect is poor and the steel cannot have good formability. Therefore, the lower limit of the Si content is made 0.01%. The Si content is made more than 0.01%. The Si content is preferably 0.05% or more, more preferably 0.10% or more, even more preferably 0.50% or more, and most preferably 0.90% or more. However, if the Si content exceeds 2.50%, it causes a decrease in the melting point of zinc, which facilitates penetration of zinc into the steel sheet during welding, resulting in a reduction in the LME resistance of the steel sheet. Therefore, the Si content is made 2.50% or less. The Si content is preferably 2.30% or less, more preferably 2.00% or less, even more preferably 1.80% or less, and most preferably 1.60% or less.
    • Mn: not less than 0.10% and not more than 5.00%
  • Mn is an element that affects the area fractions of tempered martensite, bainite, and fresh martensite through enhancement of hardenability. If the Mn content is less than 0.10%, a soft phase such as ferrite will be formed excessively, and desired area fractions of tempered martensite, bainite, and fresh martensite cannot be obtained, resulting in an insufficient steel sheet strength. Therefore, the Mn content is made 0.10% or more. The Mn content is preferably 0.50% or more, more preferably 0.80% or more, even more preferably 1.00% or more, and most preferably 2.00% or more. On the other hand, if the Mn content exceeds 5.00%, the ductility of the steel deteriorates due to increase in the area fractions of tempered martensite, bainite, and fresh martensite. Therefore, the Mn content is made 5.00% or less. The Mn content is preferably 4.50% or less, more preferably 4.00% or less, even more preferably 3.70% or less, and most preferably 3.50% or less.
    • P: 0.100% or less
  • It is possible that P, through its segregation at grain boundaries and consequential embrittlement of the grain boundaries, may adversely affect the LME resistance; therefore, the P content needs to be 0.100% or less. Thus, the P content is made 0.100% or less. The P content is preferably made 0.080% or less, more preferably 0.070% or less, even more preferably 0.050% or less, and most preferably 0.040% or less.
  • While the lower limit of the P content is not particularly limited, in view of the fact that P is a solid solution strengthening element and can increase the strength of the steel sheet, the P content is preferably made 0.001% or more, more preferably 0.003% or more, and even more preferably 0.005% or more.
    • S: 0.0200% or less
  • It is possible that S, through its segregation at grain boundaries and consequential embrittlement of the steel upon hot processing, and through the formation of sulfides, may adversely affect the LME resistance; therefore, the S content needs to be 0.0200% or less. Thus, the S content is made 0.0200% or less. The S content is preferably made 0.0180% or less, more preferably 0.0150% or less, even more preferably 0.0100% or less, and most preferably 0.0050% or less. While the lower limit of the S content is not particularly limited, in view of the restrictions of production technology, the S content is preferably 0.0001% or more, more preferably 0.0005% or more, and even more preferably 0.0010% or more.
    • Al: 0.100% or less
  • Al is an element which acts as a deoxidizer and is effective in reducing inclusions in the steel, and is preferably added in a deoxidization process. On the other hand, Al increases the transformation temperature for austenitizing, resulting in inclusion of ferrite in the microstructure. Therefore, the use of Al in an amount of more than 0.100% makes it difficult to achieve a desired TS. Thus, the Al content is made 0.100% or less. The Al content is preferably made 0.080% or less, more preferably 0.070% or less, even more preferably 0.060% or less, and most preferably 0.050% or less. While the lower limit of the Al content is not particularly limited, the Al content is preferably made 0.001% or more, more preferably 0.010% or more, and even more preferably 0.020% or more.
    • N: 0.0100% or less
  • N may adversely affect the LME resistance through the formation of coarse nitrides. If the N content exceeds 0.0100%, a large amount of coarse nitrides will be formed, leading to a significant deterioration of LME resistance. The N content is preferably as small as possible. Thus, the N content is made 0.0100% or less. The N content is preferably 0.0090% or less, more preferably 0.0080% or less, even more preferably 0.0070% or less, and most preferably 0.0060% or less. While the lower limit of the N content is not particularly limited, in view of the restrictions of production technology, the N content is preferably made 0.0001% or more, more preferably 0.0010% or more, and even more preferably 0.0020% or more.
    • O: 0.0100% or less
  • O exists as an oxide and reduces the ductility of the steel sheet. Therefore, the O content needs to be 0.0100% or less. Thus, the O content is made 0.0100% or less. The O content is preferably 0.0075% or less, more preferably 0.0060% or less, even more preferably 0.0050% or less, and most preferably 0.0045% or less. While the lower limit of the O content is not particularly specified, in view of the restrictions of production technology, the O content is preferably made 0.0001% or more, more preferably 0.0005% or more, and even more preferably 0.0010% or more.
    • Ti: not less than 0.010% and not more than 0.200%
  • Ti contributes to precipitation strengthening and, in addition, reduces the prior austenite grain size and thereby reduces the grain sizes of tempered martensite and bainite. Thus, Ti is effective in increasing the strength of the steel. In order to fully achieve this effect, the Ti content is made 0.010% or more. The Ti content is preferably 0.012% or more, more preferably 0.015% or more, even more preferably 0.020% or more, and most preferably 0.025% or more. However, if the Ti content exceeds 0.200%, Ti may remain undissolved during heating of the steel material before hot rolling, resulting in an increase in the amount of coarse precipitates and a reduction in ductility. Therefore, the Ti content is made 0.200% or less. The Ti content is preferably 0.180% or less, more preferably 0.150% or less, even more preferably 0.100% or less, and most preferably 0.050% or less.
    • Nb: not less than 0.005% and not more than 0.500%
  • Nb is an element that improves the LME resistance. In particular, the presence of dissolved Nb in the steel during welding has a great effect on the improvement of the LME resistance. More specifically, by dissolving Nb or by controlling the particle radius of Nb precipitates within a particular range, the dissolved element which is present in the steel during welding or the dissolved element produced by dissolution of the precipitates prevents the penetration of zinc into the steel, thereby improving the LME resistance. In order to fully achieve such an effect, the Nb content needs to be 0.005% or more. The Nb content is preferably 0.007% or more, more preferably 0.008% or more, even more preferably 0.010% or more, and most preferably 0.012% or more. On the other hand, if the Nb content exceeds 0.500%, precipitates such as Nb carbide and Nb nitride are formed coarsely, and the precipitates are not dissolved sufficiently during welding. Therefore, the improvement of LME resistance by the dissolved element cannot be expected. Furthermore, the El of the steel sheet will be insufficient for the TS of the steel sheet. It is quite conceivable that this may cause the occurrence of new cracking starting from coarse precipitates not only during a manufacturing process but after press forming or during welding. Therefore, the Nb content is made 0.500% or less. The Nb content is preferably made 0.400% or less, more preferably 0.350% or less, even more preferably 0.200% or less, and most preferably 0.100% or less.
  • A high-strength steel sheet according to an embodiment of the present invention has a chemical composition containing the above-described components, with the balance including Fe and incidental impurities. Preferably, a high-strength steel sheet according to an embodiment of the present invention has a chemical composition containing the above-described components, with the balance consisting of Fe and incidental impurities. Examples of the incidental impurities may include Zn, Pb, and As. Such impurities are allowed to be present in the steel in a total amount of up to 0.100%.
  • In addition to the above-described components, the chemical composition may contain, in mass %, one or two or more selected from the following: V: 0.500% or less, Ta: 0.10% or less, W: 0.10% or less, B: 0.0100% or less, Cr: 1.00% or less, Mo: 1.00% or less, Co: 1.00% or less, Ni: 1.00% or less, Cu: 1.00% or less, Sn: 0.200% or less, Sb: 0.200% or less, Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less, Zr: 0.100% or less, Te: 0.100% or less, Hf: 0.10% or less, and Bi: 0.200% or less.
    • V: 0.500% or less
  • V contributes to precipitation strengthening and, in addition, reduces the prior austenite grain size and thereby reduces the grain sizes of tempered martensite and bainite. Thus, V is effective in increasing the strength of the steel. The steel can therefore contain V as necessary. While the lower limit of the V content is not particularly limited, in order to achieve the above effect, the V content is preferably made 0.001% or more, more preferably 0.005% or more, and even more preferably 0.010% or more. However, if the V content exceeds 0.500%, V may remain undissolved during heating of the steel material before hot rolling, resulting in an increase in the amount of coarse precipitates and a reduction in ductility. Therefore, when the steel contains V, the V content is made 0.500% or less. The V content is preferably 0.400% or less, more preferably 0.300% or less, even more preferably 0.200% or less, and most preferably 0.100% or less.
    • Ta: 0.10% or less
  • Ta, like Ti, contributes to increasing the strength of the steel through the formation of an alloy carbide and an alloy carbonitride. In addition, Ta partly dissolves in Nb carbide or Nb carbonitride to form a complex precipitate such as (Nb,Ta) (C,N), thereby significantly suppressing coarsening of precipitates and stabilizing the contribution of precipitation strengthening to the increase of the strength. The steel can therefore contain Ta as necessary. While the lower limit of the Ta content is not particularly limited, in order to achieve the above effect, the Ta content is preferably made 0.01% or more, more preferably 0.02% or more, and even more preferably 0.03% or more. However, a too-high Ta content will lead to saturation of the precipitate stabilizing effect and increased alloying cost. Therefore, when the steel contains Ta, the Ta content is made 0.10% or less. The Ta content is preferably made 0.08% or less, more preferably 0.07% or less, even more preferably 0.06% or less, and most preferably 0.05% or less.
    • W: 0.10% or less
  • The steel can contain W as necessary to improve the hardenability of the steel and to further increase the strength of the steel through a reduction in the grain sizes of tempered martensite and bainite. While the lower limit of the W content is not particularly limited, in order to achieve the above effects, the W content is preferably made 0.01% or more, more preferably 0.02% or more, and even more preferably 0.03% or more. However, a W content of more than 0.10% may increase the amount of coarse precipitates, such as WN and WS, which remain undissolved during heating of the slab in a hot rolling process, resulting in a reduction in ductility. Therefore, when the steel contains W, the W content is made 0.10% or less. The W content is preferably made 0.08% or less, more preferably 0.07% or less, even more preferably 0.06% or less, and most preferably 0.05% or less.
    • B: 0.0100% or less
  • B is an element which, through its segregation at austenite grain boundaries, can improve hardenability. B can form a microstructure composed mainly of tempered martensite and bainite, thereby increasing the strength of the steel sheet, and can contribute to improving the LME resistance. Therefore, the steel can contain B as necessary. While the lower limit of the B content is not particularly limited, in order to achieve the above effects, the B content is preferably made 0.0003% or more, more preferably 0.0005% or more, and even more preferably 0.0007% or more. However, if the B content exceeds 0.0100%, coarse precipitates will be formed, resulting in a reduction in ductility. Therefore, when the steel contains B, the B content is made 0.0090% or less. The B content is preferably made 0.0080% or less, more preferably 0.0070% or less, even more preferably 0.0050% or less, and most preferably 0.0030% or less.
    • Cr: 1.00% or less
  • Cr has the effect of improving the balance between strength and ductility. Therefore, the steel can contain Cr as necessary. While the lower limit of the Cr content is not particularly limited, in order to achieve the above effect, the Cr content is preferably made 0.01% or more, more preferably 0.05% or more, and even more preferably 0.07% or more. However, if the steel contains Cr in an amount in excess of 1.00%, the area fraction of fresh martensite will be too high, resulting in a reduction in the dimensional accuracy and the ductility of the steel sheet during forming. Therefore, when the steel contains Cr, the Cr content is made 1.00% or less. The Cr content is preferably made 0.80% or less, more preferably 0.60% or less, even more preferably 0.50% or less, and most preferably 0.30% or less.
    • Mo: 1.00% or less
  • Mo has the effect of improving the balance between strength and ductility. Therefore, the steel can contain Mo as necessary. While the lower limit of the Mo content is not particularly limited, in order to achieve the above effect, the Mo content is preferably made 0.01% or more, more preferably 0.05% or more, and even more preferably 0.07% or more. However, if the steel contains Mo in an amount in excess of 1.00%, the area fraction of fresh martensite will be too high, resulting in a reduction in the dimensional accuracy and the ductility of the steel sheet during forming. Therefore, when the steel contains Mo, the Mo content is made 1.00% or less. The Mo content is preferably made 0.80% or less, more preferably 0.50% or less, even more preferably 0.30% or less, and most preferably 0.20% or less.
    • Co: 1.00% or less
  • Co is an element which is effective in improving the hardenability of the steel, and thus in strengthening the steel. Therefore, the steel can contain Co as necessary. While the lower limit of the Co content is not particularly limited, in order to achieve the above effect, the Co content is preferably made 0.01% or more, more preferably 0.05% or more, and even more preferably 0.07% or more. However, if the steel contains Co in an amount in excess of 1.00%, the area fraction of fresh martensite will be too high, resulting in a reduction in the dimensional accuracy and the ductility of the steel sheet during forming. Therefore, when the steel contains Co, the Co content is made 1.00% or less. The Co content is preferably made 0.80% or less, more preferably 0.60% or less, even more preferably 0.30% or less, and most preferably 0.20% or less.
    • Ni: 1.00% or less
  • Ni increases the strength of the steel through solid solution strengthening. Therefore, the steel can contain Ni as necessary. While the lower limit of the Ni content is not particularly limited, in order to achieve the above effect, the Ni content is preferably made 0.01% or more, more preferably 0.05% or more, and even more preferably 0.07% or more. However, if the steel contains Ni in an amount in excess of 1.00%, the area fraction of fresh martensite will be too high, resulting in a reduction in the dimensional accuracy and the ductility of the steel sheet during forming. Therefore, when the steel contains Ni, the Ni content is made 1.00% or less. The Ni content is preferably made 0.80% or less, more preferably 0.60% or less, even more preferably 0.30% or less, and most preferably 0.20% or less.
    • Cu: 1.00% or less
  • Cu is an element which is effective in strengthening the steel. Therefore, the steel can contain Cu as necessary. While the lower limit of the Cu content is not particularly limited, in order to achieve the above effect, the Cu content is preferably made 0.01% or more, more preferably 0.05% or more, and even more preferably 0.07% or more. However, if the steel contains Cu in an amount in excess of 1.00%, the area fractions of tempered martensite, bainite, and fresh martensite will be too high, resulting in a reduction in the dimensional accuracy and the ductility of the steel sheet during forming. Therefore, when the steel contains Cu, the Cu content is made 1.00% or less. The Cu content is preferably made 0.80% or less, more preferably 0.60% or less, even more preferably 0.30% or less, and most preferably 0.20% or less.
    • Sn: 0.200% or less, Sb: 0.200% or less
  • Sn and Sb suppress decarburization in a surface region having a thickness of about a few tens of µm, caused by nitridation or oxidation of the surface of the steel sheet, thereby preventing a reduction in the area fraction of tempered martensite in the steel sheet surface. In addition to this effect, Sn and Sb have the effect of ensuring the strength and the material stability. Therefore, the steel can contain these elements as necessary. While the lower limit of the content of each element is not particularly limited, in order to achieve the above effects, the content of each element is preferably made 0.001% or more, more preferably 0.003% or more, and even more preferably 0.005% or more. However, if the content of each element exceeds 0.200%, the ductility of the steel sheet may be reduced due to embrittlement of the steel sheet. Therefore, when the steel contains Sn and Sb, the content of each element is made 0.200% or less. The content of each element is preferably 0.100% or less, more preferably 0.070% or less, even more preferably 0.050% or less, and most preferably 0.030% or less.
    • Ca: 0.0100% or less, Mg: 0.0100% or less, REM: 0.0100% or less
  • The inclusion of Ca, Mg and REM in the steel, each in an amount of 0.0100% or less, causes no increase in the amount of coarse precipitates and inclusions and has no effect on the precipitation of Nb, and thus does not deteriorate the LME resistance. Therefore, when the steel contains Ca, Mg and REM, the content of each element is made 0.0100% or less. The content of each element is preferably made 0.0050% or less, more preferably 0.0040% or less, even more preferably 0.0035% or less, and most preferably 0.0030% or less. While the lower limit of the content of each element is not particularly specified, in view of the fact that each element makes the shape of a nitride or sulfide spherical and improves the limiting deformability of the steel sheet, the content of each of Ca, Mg and REM is preferably made 0.0001% or more, more preferably 0.0005% or more, even more preferably 0.0007% or more, and most preferably 0.0010% or more.
    • Zr: 0.100% or less, Te: 0.100% or less
  • The inclusion of Zr and Te in the steel, each in an amount of 0.100% or less, causes no increase in the amount of coarse precipitates and inclusions and has no effect on the precipitation of Nb, and thus does not deteriorate the LME resistance. Therefore, when the steel contains Zr and Te, the content of each element is made 0.100% or less. The content of each element is preferably made 0.080% or less, more preferably 0.070% or less, even more preferably 0.060% or less, and most preferably 0.050% or less. While the lower limit of the content of each element is not particularly specified, in view of the fact that each element makes the shape of a nitride or sulfide spherical and improves the limiting deformability of the steel sheet, the content of each of Zr and Te is preferably made 0.001% or more, more preferably 0.010% or more, and even more preferably 0.020% or more.
    • Hf: 0.10% or less
  • The inclusion of Hf in the steel in an amount of 0.10% or less causes no increase in the amount of coarse precipitates and inclusions and has no effect on the precipitation of Nb, and thus does not deteriorate the LME resistance. Therefore, when the steel contains Hf, the Hf content is made 0.10% or less. The Hf content is preferably made 0.080% or less, more preferably 0.070% or less, even more preferably 0.060% or less, and most preferably 0.050% or less. While the lower limit of the Hf content is not particularly specified, in view of the fact that Hf makes the shape of a nitride or sulfide spherical and improves the limiting deformability of the steel sheet, the Hf content is preferably made 0.003% or more, more preferably 0.010% or more, even more preferably 0.020% or more, and still more preferably 0.030% or more when the steel contains Hf.
    • Bi: 0.200% or less
  • The inclusion of Bi in the steel in an amount of 0.200% or less causes no increase in the amount of coarse precipitates and inclusions and has no effect on the precipitation of Nb, and thus does not deteriorate the LME resistance. Therefore, when the steel contains Bi, the Bi content is made 0.200% or less. The Bi content is preferably made 0.100% or less, more preferably 0.050% or less, even more preferably 0.030% or less, and most preferably 0.020% or less. While the lower limit of the Bi content is not particularly specified, in view of the fact that Bi reduces segregation, the Bi content is preferably made 0.001% or more, more preferably 0.005% or more, and even more preferably 0.010% or more.
  • When the content of any of the above-described components V, Ta, W, B, Cr, Mo, Ni, Co, Cu, Sn, Sb, Ca, Mg, REM, Zr, Te, Hf and Bi is less than the preferred lower limit, the component does not impair the effects of the present invention and is regarded as an incidental impurity.
  • (2) A description will now be given of the microstructure.
    • Area Fraction of Tempered Martensite and Bainite: not less than 40% and not more than 85%
  • Tempered martensite and bainite contribute to the strength of the steel sheet. In particular, to make the microstructure of the steel sheet composed mainly of tempered martensite and bainite is effective in maintaining a high strength. In order to fully achieve such an effect, the total area fraction of bainite and tempered martensite needs to be 40% or more. The total area fraction is preferably 42% or more, more preferably 45% or more, even more preferably 47% or more, and most preferably 50% or more. On the other hand, if the sum of the area fraction of bainite and the area fraction of tempered martensite exceeds 85%, the steel microstructure is occupied by hard phases including tempered martensite, bainite, and fresh martensite. In addition, such hard phases restrain retained austenite, and therefore the TRIP effect is not effectively exerted and the ductility is poor. Thus, it is difficult to achieve good formability. Therefore, the total area fraction of bainite and tempered martensite needs to be 85% or less. The total area fraction is preferably 83% or less, more preferably 82% or less, even more preferably 81% or less, and most preferably 80% or less.
    • Area Fraction of Fresh Martensite: not less than 0% and not more than 25%
  • Fresh martensite is a very hard phase, and therefore increases the strength of the steel. Fresh martensite is not necessarily required if the strength of the steel sheet is secured; however, the inclusion of fresh martensite in the steel sheet microstructure can further increase the strength of the steel sheet, achieving further strengthening. Therefore, the area fraction of fresh martensite needs to be 0% or more. The area fraction is preferably 2% or more, more preferably 3% or more, even more preferably 4% or more, and most preferably 5% or more. On the other hand, fresh martensite reduces the ductility of the steel, making it difficult to achieve good formability. Therefore, the area fraction of fresh martensite needs to be 25% or less. The area fraction is preferably 23% or less, more preferably 22% or less, even more preferably 21% or less, and most preferably 20% or less.
    • Area Fraction of Retained Austenite: not less than 5% and not more than 20%
  • Retained austenite transforms into martensite by the TRIP effect during processing. This increases the strength of the steel and, at the same time, enhances the strain dispersibility, thereby improving the ductility. In order to ensure good formability, the area fraction of retained austenite needs to be 5% or more. The area fraction is preferably 7% or more, more preferably 8% or more, even more preferably 9% or more, and most preferably 10% or more. On the other hand, if the area fraction of retained austenite exceeds 20%, elements concentrated in the retained austenite, especially Si, will diffuse during welding, causing a decrease in the melting point of zinc. This may facilitate the penetration of zinc into the steel sheet during welding, resulting in a deterioration of the LME resistance of the steel sheet. Therefore, in order to ensure sufficient ductility while preventing deterioration of the LME resistance, the area fraction of retained austenite needs to be 20% or less. The area fraction is preferably 18% or less, more preferably 17% or less, even more preferably 16% or less, and most preferably 15% or less.
  • In the present invention, the inclusion of a remaining microstructure, such as ferrite or pearlite, in the steel does not impair the effects of the present invention. When there is a remaining microstructure, it may include at least one of ferrite and pearlite, i.e., the area fraction of at least one of ferrite and pearlite may be 0% or more. The area fraction of the remaining microstructure is preferably 20% or less, more preferably 18% or less, even more preferably 15% or less, and most preferably 13% or less.
  • The relationship between the amount of dissolved Nb in the steel sheet (Nbsol), the amount of Nb in Nb precipitates having a particle size of less than 20 nm (Nbpre), and the total amount of Nb contained in the steel sheet (Nb) satisfies the following Formula 1: Nb sol / Nb + Nb pre / Nb ≥ 0.40 where Nbsol represents the amount (mass %) of dissolved Nb, and Nbpre represents the amount (mass %) of Nb in Nb precipitates having a particle size of less than 20 nm.
  • To determine the relationship between the amount of dissolved Nb in the steel sheet (Nbsol, unit: mass %), the amount of Nb in Nb precipitates having a particle size of less than 20 nm (Nbpre, unit: mass %), and the total amount of Nb contained in the steel sheet (Nb, unit: mass %) is one of the important factors of the present invention. In the present invention, Nb in the steel sheet is brought to a dissolved state, or to a state which can be converted into a dissolved state during welding, thereby preventing non-self-derived LME and improving the LME resistance. The present inventors, through experiments conducted using various steel sheets with varying states of Nb present therein, have found that the LME resistance is improved when Formula 1 is satisfied.
  • The mechanism by which Nb in the steel being in a dissolved state during welding improves the LME resistance is not clear, it is conceivable that when Nb exists in a dissolved state, it prevents the penetration of zinc into the steel sheet, which is the direct cause of LME cracking. Formula 1 thus needs to be satisfied. "(Nbsol/Nb) + (Nbpre/Nb)" is preferably 0.41 or more, more preferably 0.42 or more. While the upper limit is not particularly limited, "(Nbsol/Nb) + (Nbpre/Nb)" is preferably 0.95 or less, more preferably 0.90 or less. The lower limit of the particle size of the Nb precipitates is not particularly limited; for example, the Nb precipitates may have a particle size of 0.1 nm or more.
    • Amount of Diffusible Hydrogen in Steel: 0.50 mass ppm or less
  • In order to ensure good ductility and thus good formability, the amount of diffusible hydrogen in the steel is made 0.50 mass ppm or less. The amount of diffusible hydrogen is preferably made 0.30 mass ppm or less, more preferably 0.25 mass ppm or less, even more preferably 0.20 mass ppm or less, and most preferably 0.15 mass ppm or less. While the lower limit of the amount of diffusible hydrogen in the steel is not particularly specified, in view of the restrictions of production technology, the amount of diffusible hydrogen is preferably made 0.01 mass ppm or more, more preferably 0.02 mass ppm or more, even more preferably set to 0.03 mass ppm or more, and most preferably 0.05 mass ppm or more.
  • (3) A manufacturing method and manufacturing conditions will now be described.
    • [Slab (Steel Material) Heating Step]
  • Steel Material Heating Conditions: heating for 1.0 hours or more at a temperature equal to or higher than Tsol°C represented by Formula 2
    Precipitates existing in the heating stage of the steel material will exist as coarse precipitates in the steel sheet finally obtained, which not only adversely affects the LME resistance, but does not contribute to the strength, and thus does not improve the TS × El balance. Therefore, coarse precipitates, which have been precipitated during casting, need to be re-dissolved as much as possible. If the steel material heating temperature is below Tsol°C or the heating time is less than one hour, the dissolution of Nb will be insufficient, and Formula 1 will not be satisfied. This may result in poor LME resistance and in an increased risk of trouble during hot rolling due to an increase in rolling load. Also from the viewpoint of scaling off defects such as bubbles and segregations in the surface of the steel material, and reducing cracking and unevenness on the surface of the steel sheet, thereby achieving a smooth surface, the steel material heating temperature needs to be equal to or higher than Tsol°C. Therefore, the steel material heating temperature is made equal to or higher than Tsol°C represented by the following Formula 2. The steel material heating temperature is preferably equal to or higher than "Tsol × 1.1"°C, more preferably equal to or higher than "Tsol × 1.2"°C, even more preferably equal to or higher than "Tsol × 1.3"°C, and most preferably equal to or higher than "Tsol × 1.5"°C. The heating time at the temperature is made 1.0 hours or more. The heating time is preferably 1.1 hours or more, more preferably 1.2 hours or more, even more preferably 1.3 hours or more, and most preferably 1.5 hours or more. While the upper limit of the steel material heating temperature is not particularly specified, the use of a heating temperature exceeding 1500°C increases the amount of oxidation, resulting in an increased scale loss. Therefore, the steel material heating temperature is preferably made 1500°C or lower, more preferably 1450°C or lower, and even more preferably 1400°C or lower. T sol = 7900 / 3.42 − log Nb % C % − 273
    • [Hot Rolling Step] Finish Rolling Start Temperature in Hot Rolling: TFET°C
  • The steel material after the heating is hot-rolled into a hot-rolled steel sheet. While the upper and lower limits of the start temperature of finish rolling in the hot rolling are not particularly specified, a finish rolling start temperature of less than "Tsol°C - 100°C" will cause precipitation of Nb which has been dissolved during heating of the steel material, resulting in an increase in the particle size of precipitates. Accordingly, Formula 1 may not be satisfied, and thus the LME resistance may be poor. Therefore, the finish rolling start temperature is preferably made "Tsol°C - 100°C" or higher, more preferably "Tsol°C - 70°C" or higher, even more preferably "Tsol°C - 50°C" or higher, and most preferably "Tsol°C - 30°C" or higher. On the other hand, a finish rolling start temperature exceeding 1200°C may increase a scale loss during pre-heating of the slab, causing breakage of the sheet during hot rolling. Therefore, the finish rolling start temperature in hot rolling is preferably made 1200°C or lower, more preferably 1170°C or lower, even more preferably 1150°C or lower, and most preferably 1120°C or lower.
    • Finish Rolling Delivery Temperature in Hot Rolling: TFDT°C
  • The steel material after the heating is hot-rolled into a hot-rolled steel sheet. While the upper and lower limits of the delivery temperature in the finish rolling are not particularly specified, a finish rolling delivery temperature of less than 800°C will cause precipitation of Nb which has been dissolved during heating of the steel material, resulting in an increase in the amount of Nb precipitated. Further, such a delivery temperature increases a rolling load, and thus a rolling burden, which may impede a cold rolling process. Therefore, the finish rolling delivery temperature in hot rolling is preferably made 800°C or higher, more preferably 820°C or higher, even more preferably 840°C or higher, and most preferably 850°C or higher. If the finish temperature exceeds 1000°C, the amount of oxides (scale) produced tends to increase rapidly, resulting in roughening of the interface between the base steel and the oxides, and deterioration of the surface quality after pickling and cold rolling. In addition, crystal grains may become excessively coarse, causing surface roughening of the pressed product during processing. Therefore, the finish rolling delivery temperature is preferably made 1000°C or lower, more preferably 980°C or lower, even more preferably 970°C or lower, and most preferably 950°C or lower.
    • Effective Time: tHR
  • Effective time tHR is herein defined as the time required from the start of finish rolling to the end of finish rolling. Since Nb contained in the steel sheet begins to precipitate and grow during finish rolling, the effective time tHR is one of the parameters for controlling the form of Nb to improve the LME resistance. While the upper and lower limits of the effective time tHR are not particularly specified, the effective time tHR is preferably made 3 seconds or more, more preferably 4 seconds or more, even more preferably 5 seconds or more, and most preferably 7 seconds or more. The effective time tHR is preferably made 15 seconds or less, more preferably 12 seconds or less, even more preferably 11 seconds or less, and most preferably 10 seconds or less.
    • [Coiling Step] Residence Time from TFDT°C to 650°C: tCT
  • Residence time tCT is herein defined as the time from the finish rolling delivery temperature to 650°C. Since the precipitation of Nb and the growth of precipitated Nb are likely to proceed during the period from the completion of finish rolling to the sheet temperature reaching 650°C, the residence time tCT is one of the parameters for controlling the form of Nb to improve the LME resistance. While the upper and lower limits of the residence time tCT are not particularly specified, the residence time tCT is preferably made 5 seconds or more, more preferably 7 seconds or more, even more preferably 8 seconds or more, and most preferably 9 seconds or more. The residence time tCT is made 20 seconds or less, more preferably 18 seconds or less, even more preferably 17 seconds or less, and most preferably 15 seconds or less.
    • Coiling Temperature TCT°C after Hot Rolling: 650°C or lower
  • If the coiling temperature after hot rolling is higher than 650°C, the precipitation of Nb and the growth of precipitated Nb may proceed excessively, so that Formula 1 may not be satisfied. This causes deterioration of the LME resistance and, in addition, causes the formation of an oxide film, which is difficult to remove by pickling, on the surface of the hot-rolled sheet, which may impair the surface appearance of the steel sheet after cold rolling. Therefore, the coiling temperature after hot rolling is made 650°C or lower. The coiling temperature is preferably 630°C or lower, more preferably 620°C or lower, even more preferably 610°C or lower, and most preferably 600°C or lower. While the lower limit of the coiling temperature is not particularly limited, if the coiling temperature is lower than 300°C, the strength of the hot-rolled sheet increases, which may cause an increase in the rolling burden in cold rolling and a defective sheet shape, resulting in a reduction in productivity. Therefore, the lower limit of the coiling temperature is preferably made 300°C or higher, more preferably 320°C or higher, even more preferably 350°C or higher, and most preferably 400°C or higher.
  • The resulting hot-rolled steel sheet (hot-rolled coil) may be subjected to intermediate heat treatment at a temperature below 650°C, as necessary, to prevent an increase in load in later cold rolling. The heat treatment is preferably performed at a temperature of 150°C or higher. The heat treatment may be carried out, for example, in a box annealing furnace at a soaking temperature of 500°C for a soaking time of 4 hours.
  • The resulting hot-rolled steel sheet (hot-rolled coil) may be subjected to a treatment, such as pickling, as necessary. Pickling of the hot-rolled coil may be performed by a common method. The hot-rolled coil may also be subjected to skin pass rolling to correct its shape and enhance its pickling properties.
  • After the hot rolling and/or the intermediate heat treatment and/or the pickling, the steel sheet may be subjected directly to the below-described annealing step (heat treatment), or may be subjected to cold rolling prior to the heat treatment. When cold rolling is performed, the cold rolling reduction is preferably made 25% or more, more preferably 30% or more, even more preferably 32% or more, and most preferably 35% or more. On the other hand, an excessive reduction causes a too-high rolling load, leading to an increase in the burden on a cold rolling mill. Therefore, the cold rolling reduction is preferably made 75% or less, more preferably 70% or less, even more preferably 67% or less, and most preferably 65% or less.
    • [Annealing Step] Heating Time from 650°C to Soaking Temperature TAT°C: t
  • The precipitation of Nb and the growth of precipitated Nb proceeds at a sheet temperature of 650°C or higher. Thus, the heating time t is one of the parameters for controlling the form of Nb to improve the LME resistance. While the upper and lower limits of the heating time t are not particularly specified, the heating time t is preferably made 300 seconds or more, more preferably 400 seconds or more, even more preferably 450 seconds or more, and most preferably 490 seconds or more. The heating time t is preferably made 700 seconds or less, more preferably 650 seconds or less, even more preferably 600 seconds or less, and most preferably 590 seconds or less.
    • Soaking Temperature: TAT not less than 750°C and not more than 950°C
  • When the steel sheet is held in a temperature range below 750°C, it is held in a two-phase region. In such a case, the total area fraction of tempered martensite and bainite in the final microstructure is less than 40%, and therefore a sufficient steel sheet strength cannot be ensured. Therefore, the soaking temperature TAT°C is made 750°C or higher, preferably 770°C or higher, more preferably 800°C or higher, and even more preferably 840°C or higher. On the other hand, if the steel sheet is held in a temperature range above 950°C, the precipitation of Nb is promoted, and the growth of Nb precipitates proceeds excessively. Accordingly, Formula 1 will not be satisfied, leading to a reduction in the LME resistance. Therefore, the soaking temperature TAT°C is made 950°C or lower, preferably 940°C or lower, more preferably 930°C or lower, and even more preferably 920°C or lower.
    • Holding Time at Soaking Temperature TAT°C: tAT
  • Holding time tAT at the soaking temperature TAT°C is one of the control parameters for promoting austenitizing of the steel sheet during soaking and, since it is related to the precipitation of Nb and the growth of Nb precipitates, is also one of the parameters for controlling the LME resistance. While the upper and lower limits of the holding time tAT are not particularly specified, the holding time tAT is preferably made 15 seconds or more in order to sufficiently promote austenitizing. The holding time tAT is more preferably made 30 seconds or more, even more preferably 50 seconds or more, and most preferably 100 seconds or more. On the other hand, from the viewpoint of making the particles of Nb precipitates finer to improve the LME resistance, the holding time tAT is preferably made 1000 seconds or less, more preferably 500 seconds or less, even more preferably 400 seconds or less, and most preferably 350 seconds or less.
    • Sum of QHR, QCT and QAT: 6000 or less
  • In order to suppress the precipitation of Nb and the growth of Nb precipitates during heating in the hot rolling step, the coiling step, and the annealing step, the following is an important requirement in the present invention for reasons which will be described below: QHR defined by the following Formula 3 in terms of the finish rolling start temperature TFET°C, the finish rolling delivery temperature TFDT°C, and the effective time tHR from the start of finish rolling to the end of finish rolling, QCT defined by the following Formula 4 in terms of the residence time tCT from TFDT°C to 650°C and the coiling temperature TCT°C in the coiling step, and QAT defined by the following Formula 5 in terms of the soaking temperature TAT°C, the heating time t from 650°C to the soaking temperature TAT°C, and the holding time tAT at the soaking temperature TAT°C in the annealing step, satisfy the following Formula 6, i.e., the sum of QHR, QCT, and QAT is 6000 or less. Q HR = 0.5 T FET + T FDT × log 10 t HR Q CT = 0.5 T CT + 650 × log 10 t CT Q AT = 0.5 650 + T AT × log 10 t + T AT × log 10 t AT Q HR + Q CT + Q AT ≤ 6000
  • That the sum of QHR, QCT, and QAT is 6000 or less is an important requirement in the present invention in order to suppress the precipitation and growth of Nb in the hot rolling step, the coiling step, and the heating step. As described above, QHR, QCT, and QAT are each defined by the temperatures at the start and end of each of the effective time tHR, the residence time tCT, the heating time t, and the holding time tAT, during which heat of 650°C to Tsol is applied.
  • If the sum of QHR, QCT, and QAT exceeds 6000, Nb will be precipitated excessively and the precipitates will grow excessively, and thus the amount of Nb that improves the LME resistance will decrease, resulting in poor LME resistance. Therefore, the sum of QHR, QCT, and QAT is made 6000 or less, preferably 5990 or less, more preferably 5980 or less, even more preferably 5970 or less, and most preferably 5950 or less. While the lower limit of the sum of QHR, QCT, and QAT is not particularly specified, in view of the feasible range in actual operations, the sum is preferably made 3500 or more, more preferably 3600 or more, even more preferably 3700 or more, most preferably 3800 or more, and particularly preferably 3900 or more.
    • QAT Value of 4700 or Less (Preferred Condition)
  • From the viewpoint of taking account of the total heat applied during heating in the hot rolling step, the coiling step, and the annealing step to suppress excessive precipitation and growth of Nb, its evaluation is performed based on the sum of QHR, QCT, and QAT. Therefore, the upper and lower limits of each parameter are not particularly specified. However, from the viewpoint of suppressing excessive growth of Nb during heating in the annealing step, which directly leads to the final microstructure of the steel sheet, QAT is preferably made 4700 or less, more preferably 4650 or less, even more preferably 4630 or less, and most preferably 4600 or less. From the viewpoint of more precisely controlling the austenite fraction during annealing in order to increase the strength of the steel, QAT is preferably made 3500 or more, more preferably 3800 or more, even more preferably 4100 or more, and most preferably 4120 or more.
  • Likewise for the hot rolling step and the coiling step, from the viewpoint of suppressing excessive growth of Nb, QHR is preferably made 1200 or less, more preferably 1170 or less, and even more preferably 1150 or less. QCT is preferably made 850 or less, more preferably 840 or less, and even more preferably 835 or less. On the other hand, while no particular limitation is placed especially on the lower limit of QHR, QHR is preferably made 250 or more, more preferably 600 or more, and even more preferably 700 or more. QCT is preferably made 150 or more, more preferably 170 or more, and even more preferably 180 or more.
  • In the annealing step, after the steel sheet is held at the soaking temperature TAT for the holding time tAT, the steel sheet may be simply cooled, may be cooled to an arbitrary holding temperature and held at the temperature, and then cooled again, or may be cooled to an arbitrary cooling stop temperature, then reheated to a holding temperature and held at the temperature, and then cooled again.
  • After the annealing step, the steel sheet may be reheated. When martensite is formed in the steel sheet and then tempered, C is concentrated in untransformed austenite. This enhances the stability of austenite and increases the area fraction of retained austenite contained in the steel sheet after cooling, making it possible to further enhance the ductility. When the steel sheet is reheated for such reasons, it is preferred to cool the steel sheet to a cooling stop temperature of not less than 100°C and not more than 350°C, reheat the steel sheet to a holding temperature of not less than 200°C and not more than 450°C, hold the steel sheet at the temperature, and then cool the steel sheet to room temperature. The reheating temperature is preferably made 200°C or higher, more preferably 210°C or higher, even more preferably 230°C or higher, and most preferably 250°C or higher. The reheating temperature is preferably made 450°C or lower, more preferably 430°C or lower, even more preferably 410°C or lower, and most preferably 400°C or lower.
  • When reheating of the steel sheet is not performed, it is preferred to hold the steel sheet at a holding temperature of not less than 200°C and not more than 450°C for 10 seconds or more, and then cool the steel sheet to room temperature. The steel sheet is more preferably held at a holding temperature of not less than 200°C and not more than 450°C for 20 seconds or more, even more preferably for 25 seconds or more. While the upper limit of the holding time is not particularly limited, the holding time is preferably made 500 seconds or less, more preferably 400 seconds or less.
    • Coating Treatment
  • When the steel sheet is subjected to hot-dip galvanizing, for example, the steel sheet which has undergone the above-described annealing treatment is immersed in a galvanizing bath at 440°C to 500°C to perform hot-dip galvanizing of the steel sheet, followed by gas wiping or the like to adjust the coating weight. The temperature of the galvanizing bath is preferably made 440°C or higher, more preferably 450°C or higher, and even more preferably 455°C or higher. The temperature of the galvanizing bath is preferably made 500°C or lower, more preferably 490°C or lower, and even more preferably 485°C or lower.
  • While there is no particular restriction placed on the coating conditions, the coating weight (per one surface) is preferably 20 g/m2 or more from the viewpoint of corrosion resistance and control of the coating weight. The coating weight is more preferably 25 g/m2 or more, even more preferably 30 g/m2 or more, and most preferably 32 g/m2 or more. From the viewpoint of adhesion, the coating weight is preferably 120 g/m2 or less, more preferably 100 g/m2 or less, even more preferably 70 g/m2 or less, and most preferably 65 g/m2 or less.
  • The hot-dip galvanizing is preferably performed using a galvanizing bath containing 0.08% to 0.30% of Al. Thus, the amount of Al in the hot-dip galvanizing is preferably made 0.08% or more, more preferably 0.09% or more, even more preferably 0.10% or more, and most preferably 0.12% or more. The amount of Al in the hot-dip galvanizing is preferably made 0.30% or less, more preferably 0.25% or less, even more preferably 0.22% or less, and most preferably 0.20% or less. If elements other than Al, Mg, and Si, such as Pb, Sb, Fe, Mg, Mn, Ni, Ca, Ti, V, Cr, Co, and Sn are mixed into the galvanizing bath, such elements will not significantly affect the effects of the present invention.
  • When an alloying treatment is performed on a galvanized coating, the treatment is performed in a temperature range of 450°C to 600°C. If the alloying treatment is performed at a temperature exceeding 600°C, untransformed austenite will be transformed into pearlite, and the area fraction of retained austenite will be less than 5%, which may result in a reduction in ductility. Therefore, when an alloying treatment of a galvanized coating is performed, the alloying treatment is preferably performed in a temperature range of not less than 450°C, more preferably not less than 460°C, even more preferably not less than 465°C, and most preferable not less than 470°C. Further, when an alloying treatment of a galvanized coating is performed, the alloying treatment is preferably performed in a temperature range of not more than 600°C, more preferably not more than 570°C, even more preferably not more than 550°C, and most preferably not more than 530°C. The alloyed coating layer of the hot-dip galvanized steel sheet preferably has an Fe concentration of 8% to 17%. Thus, the Fe concentration of the alloyed coating layer of the hot-dip galvanized steel sheet is preferably 8% or more, more preferably 9% or more, and even more preferably 10% or more. The Fe concentration of the alloyed coating layer of the hot-dip galvanized steel sheet is preferably 17% or less, more preferably 16% or less, and even more preferably 15% or less. The alloyed coating layer is thus formed by performing the alloying treatment on the hot-dip galvanized steel sheet.
    • EXAMPLES
  • Molten steels having the chemical compositions shown in Tables 1-1 and 1-2, with the balance being Fe and incidental impurities, were produced in a converter. Each molten steel was continuously cast into a slab. The slab was subjected to a heating step, a hot rolling step, a coiling step, and an annealing (heating) step under the conditions shown in Tables 2-1 and 2-2. When reheating of the steel sheet was performed, the steel sheet after annealing was cooled to a cooling stop temperature of 100°C to 350°C, reheated to a holding temperature of 200°C to 450°C, held at the holding temperature, and then cooled to room temperature. When reheating of the steel sheet was not performed, the steel sheet after annealing was held at a holding temperature of 200°C to 450°C for 10 seconds or more, and then cooled to room temperature. In either case, the cooled steel sheet was cold-rolled at a rolling reduction of 50% to obtain a high-strength cold-rolled steel sheet (CR). Some steel sheets after the coiling step were subjected to the heating step, but were not subjected to cold rolling, to obtain hot-rolled steel sheets (HR). Some steel sheets were subjected to hot-dip galvanizing to obtain hot-dip galvanized steel sheets (GI) and galvannealed steel sheets (GA). A galvanizing bath containing 0.19 mass % Al was used for the hot-dip galvanized steel sheets (GI), and a galvanizing bath containing 0.14 mass % Al was used for the galvannealed steel sheets (GA). The temperature of each bath was 465°C. The coating weight was 45 g/m2 per one surface (double-sided coating). For the galvannealed steel sheets (GA), the Fe concentration in the coating layer was adjusted to be within the range of 9 mass % to 12 mass %.
  • Each of the resulting steel sheets was examined for its cross-sectional microstructure, tensile properties, and LME resistance, the results of which are shown in Tables 3-1 and 3-2. [Table 2-1]
    Heating step Hot rolling step Coiling step Annealing step
    Steel sheet sample No. Steel symbol Tsol (°C) Heating temp. T (°C) Heating time (hour) Finish rolling start temp. TFET (°C) Finish rolling delivery temp. TFDT (°C) Effective time tHR (sec) QHR (°C·sec) Residence time tCT (sec) Coiling temp. TCT (°C) QCT (°C·sec) Heating time t (sec) Soaking temp. TAT (°C) Holding time tAT (sec) QAT (°C·sec) QHR + QCT + QAT (°C·sec) Re- heated? State of steel sheet
    1 A 899 1270 1.8 1080 920 6 778 14 560 693 580 880 280 4268 5739 yes CR
    2 A 899 1270 1.8 1080 920 7 845 6 510 451 480 870 260 4139 5435 no CR
    3 B 1038 1260 1.3 1110 970 7 879 12 574 660 570 920 260 4385 5924 yes CR
    4 B 1038 1260 1.3 1060 970 4 611 11 590 646 560 880 240 4197 5454 no CR
    5 C 1072 1290 1.1 1060 940 6 778 11 590 646 570 850 280 4147 5571 yes CR
    6 C 1072 1290 1.1 1100 940 8 921 15 500 676 550 860 230 4100 5697 no CR
    7 C 1072 1290 1.1 1070 940 6 782 11 490 594 540 920 250 4351 5727 yes GA
    8 C 1072 1290 1.1 1110 940 6 798 9 540 568 580 920 290 4435 5801 yes GI
    9 C 1072 1250 1.4 1110 950 6 801 8 550 542 580 900 290 4358 5701 yes HR
    10 D 1101 1280 1.3 1110 900 7 849 6 570 475 500 890 290 4270 5594 yes CR
    11 D 1101 1280 1.3 1110 900 7 849 9 510 553 540 860 280 4168 5570 no CR
    12 D 1101 1070 1.3 1040 900 7 820 13 610 702 550 880 300 4276 5798 yes CR
    13 D 1101 1280 0.4 1120 900 7 854 12 540 642 580 860 280 4191 5687 yes CR
    14 D 1101 1280 1.3 1060 900 6 763 9 700 644 440 900 240 4191 5598 yes CR
    15 D 1101 1280 1.3 1060 900 4 590 13 560 674 550 980 250 4583 5847 yes CR
    16 D 1101 1280 1.3 1090 900 5 695 8 560 546 500 720 240 3563 4804 yes CR
    17 D 1101 1280 1.3 1160 960 14 1215 6 570 475 560 890 300 4321 6011 yes CR
    18 D 1101 1280 1.3 1160 960 7 896 21 640 853 560 890 300 4321 6070 yes CR
    19 D 1101 1280 1.3 1160 960 7 896 6 570 475 560 940 500 4722 6093 yes CR
    20 D 1101 1280 1.3 1160 960 7 896 15 500 676 550 950 230 4436 6008 yes CR
    21 D 1101 1280 1.3 1110 870 2 298 9 530 563 510 910 240 4278 5139 no CR
    22 E 1205 1300 1.8 1150 950 6 817 10 520 585 520 900 300 4334 5736 yes CR
    23 F 1015 1240 1.5 1120 900 5 706 14 600 716 520 870 260 4165 5587 yes CR
    24 G 1205 1300 1.3 1120 880 12 1079 8 610 569 490 920 250 4318 5966 yes CR
    25 H 1107 1140 1.4 1080 890 6 766 8 510 524 490 850 280 4098 5388 yes CR
    26 I 1013 1230 1.5 1060 970 7 858 7 510 490 520 890 120 3942 5290 yes CR
    27 J 996 1220 1.0 1060 980 6 794 8 610 569 510 920 260 4347 5710 yes CR
    28 K 1015 1470 1.9 1190 800 5 695 13 560 674 550 930 250 4395 5764 yes CR
    29 L 1140 1280 1.8 1140 860 5 699 8 580 555 480 880 240 4146 5400 yes CR
    30 M 1146 1250 1.5 1120 930 7 866 10 650 650 550 850 270 4122 5638 yes CR
    31 N 1067 1310 1.7 1130 880 6 782 9 550 573 470 870 260 4132 5487 yes CR
    32 O 785 1260 1.1 1110 950 5 720 9 500 549 440 880 240 4117 5386 yes CR
    33 P 1192 1300 1.4 1070 910 7 837 15 550 706 460 860 300 4141 5684 yes CR
    34 Q 1096 1280 1.1 1120 900 6 786 13 560 674 560 860 120 3863 5323 yes CR
    35 R 1092 1280 1.3 1070 940 7 849 12 500 621 580 920 250 4375 5845 yes CR
    36 S 1045 1270 1.2 1080 960 8 921 10 580 615 480 870 200 4040 5576 yes CR
    37 T 1076 1330 1.9 1080 910 7 841 6 580 479 520 940 260 4429 5749 yes CR
    38 U 1061 1290 1.4 1070 900 9 940 5 540 416 560 940 300 4513 5869 yes CR
    39 V 1018 1240 1.0 1100 900 4 602 10 520 585 520 900 300 4334 5521 yes CR
    40 W 1049 1230 1.0 1090 920 8 908 9 580 587 560 870 260 4190 5685 yes CR
    41 X 1170 1230 1.0 1120 900 8 912 10 500 575 570 930 250 4407 5894 yes CR
    42 Y - 1280 1.1 1090 870 9 935 12 500 621 500 950 230 4403 5959 yes CR
    43 Z 1653 1680 1.7 1090 950 5 713 10 510 580 460 920 280 4342 5635 yes CR
    * HR: hot-rolled steel sheet (no coating), CR: cold-rolled steel sheet (no coating), GI: hot-dip galvanized steel sheet (no alloying of galvanized coating), GA: galvannealed steel sheet The underlines indicate that the values are outside the scope of the present invention.
    [Table 2-2]
    Heating step Hot rolling step Coiling step Annealing step
    Steel sheet sample No. Steel symbol Tsol (°C) Heating temp. T (°C) Heating time (hour) Finish rolling start temp. TFET (°C) Finish rolling delivery temp. TFDT (°C) Effective time tHR (sec) QHR (°C·sec) Residence time tCT (sec) Coiling temp. TCT (°C) QCT (°C·sec) Heating time t (sec) Soaking temp. TAT (°C) Holding time tAT (sec) QAT (°C·sec) QHR + QCT + QAT (°C·sec) Reheated? State of steel sheet
    44 AA 1042 1320 1.4 1060 890 7 824 9 510 553 540 860 280 4168 5545 yes CR
    45 AB 1072 1310 2.0 1120 890 5 702 12 540 642 490 860 280 4136 5480 yes CR
    46 AC 1126 1330 2.0 1070 920 10 995 14 640 739 520 860 300 4181 5915 yes CR
    47 AD 1078 1290 1.6 1120 950 9 988 15 580 723 490 900 240 4227 5938 yes CR
    48 AE 1142 1270 1.1 1100 880 8 894 11 590 646 560 880 240 4197 5737 yes CR
    49 AF 1108 1310 1.3 1060 860 6 747 15 580 723 430 880 300 4194 5664 yes CR
    50 AG 1012 1280 1.1 1120 960 8 939 14 600 716 520 870 260 4165 5820 yes CR
    51 AH 1097 1230 1.0 1090 910 4 602 13 510 646 580 910 270 4368 5616 yes CR
    52 AI 1095 1310 1.7 1080 890 9 940 8 540 537 500 910 300 4359 5836 yes CR
    53 AJ 1083 1260 1.9 1050 970 6 786 16 500 692 500 920 240 4308 5786 yes CR
    54 AK 1104 1280 1.1 1070 960 6 790 10 580 615 580 890 320 4357 5762 yes CR
    55 AL 1128 1310 1.3 1080 860 7 820 14 650 745 440 910 240 4228 5793 yes CR
    56 AM 1065 1220 2.0 1120 950 9 988 9 530 563 510 910 240 4278 5829 yes CR
    57 AN 1079 1310 1.7 1050 950 6 778 11 590 646 570 850 280 4147 5571 yes CR
    58 AO 1135 1260 1.5 1110 970 7 879 12 540 642 580 920 300 4448 5969 yes CR
    59 AP 1125 1280 1.6 1060 860 6 747 13 610 702 550 880 300 4276 5725 yes CR
    60 AQ 1121 1240 1.3 1050 890 9 926 19 580 786 620 860 300 4239 5951 yes CR
    61 AR 1134 1250 1.1 1090 890 5 692 19 580 786 620 860 300 4239 5717 yes CR
    62 AS 1120 1210 1.4 1120 950 6 805 13 510 646 590 880 300 4300 5751 yes CR
    63 AT 1101 1230 1.3 1190 930 9 1011 14 530 676 510 890 250 4219 5906 yes CR
    64 AU 1100 1360 1.2 1170 990 9 1031 11 570 635 440 900 290 4265 5931 yes CR
    65 AV 1129 1250 1.6 1160 920 6 809 15 620 747 550 890 270 4274 5830 yes CR
    66 AW 1119 1290 1.8 1140 960 8 948 15 510 682 560 890 310 4333 5963 yes CR
    67 AX 1575 1580 2.0 1190 910 7 887 15 540 700 550 870 260 4184 5771 yes CR
    68 AY 1124 1240 1.8 1100 910 7 849 16 570 735 450 900 250 4214 5798 yes CR
    69 AZ 1043 1230 1.9 1100 980 5 727 10 600 625 600 910 280 4394 5746 yes CR
    70 E 1205 1290 1.7 1140 930 6 805 9 520 558 560 920 290 4423 5786 no GA
    * HR: hot-rolled steel sheet (no coating), CR: cold-rolled steel sheet (no coating), GI: hot-dip galvanized steel sheet (no alloying of galvanized coating), GA: galvannealed steel sheet The underlines indicate that the values are outside the scope of the present invention.
    [Table 3-1]
    Steel sheet microstructure Steel sheet properties LME evaluation test
    Steel sheet sample No. Steel symbol TM + B (%) FM (%) γR (%) Balance (%) Nbsol (mass %) Nbpre (mass %) Formula1 Amount of diffusible hydrogen (mass ppm) TS (MPa) EI (%) TS1.5 × EI (MPa·%) Electrode inclination angle (°) Holding time (s) Evaluation results Remarks
    1 A 72 1 14 13 0.002 0.004 0.50 0.02 1173 15.4 618682 3 0.12 ⊚ Inventive Ex.
    2 A 64 6 11 19 0.002 0.005 0.58 0.02 1180 12.7 514786 3 0.14 ⊚ Inventive Ex.
    3 B 84 4 9 3 0.002 0.007 0.45 0.03 1415 10.5 558887 5 0.08 △ Inventive Ex.
    4 B 58 16 10 16 0.003 0.008 0.55 0.02 1339 12.0 587966 3 0.14 ⊚ Inventive Ex.
    5 C 52 17 11 20 0.002 0.007 0.47 0.04 1453 8.5 470779 3 0.10 ⊚ Inventive Ex.
    6 C 59 14 9 18 0.002 0.007 0.47 0.01 1414 11.0 584880 4 0.10 ○ Inventive Ex.
    7 C 76 15 9 0 0.002 0.007 0.47 0.14 1526 10.6 631884 5 0.14 ○ Inventive Ex.
    8 C 75 16 7 2 0.002 0.007 0.47 0.18 1512 8.2 482105 3 0.08 △ Inventive Ex.
    9 C 73 18 8 1 0.002 0.007 0.47 0.11 1498 8.9 516010 4 0.10 △ Inventive Ex.
    10 D 81 6 12 1 0.003 0.007 0.48 0.03 1563 12.1 747694 3 0.08 ⊚ Inventive Ex.
    11 D 66 19 9 6 0.003 0.008 0.52 0.03 1584 9.5 598903 5 0.12 ⊚ Inventive Ex.
    12 D 71 21 8 0 0.001 0.004 0.24 0.01 1540 8.2 495559 5 0.14 × Comp. Ex.
    13 D 65 18 11 6 0.002 0.006 0.38 0.02 1476 12.4 703155 5 0.14 × Comp. Ex.
    14 D 72 18 10 0 0.002 0.005 0.33 0.02 1571 10.0 622679 3 0.14 × Comp. Ex.
    15 D 71 19 7 3 0.002 0.005 0.33 0.02 1460 7.2 401663 5 0.08 × Comp. Ex.
    16 D 25 1 13 61 0.003 0.009 0.57 0.04 953 15.4 453065 3 0.10 ⊚ Comp. Ex.
    17 D 84 3 7 6 0.002 0.005 0.33 0.03 1530 7.8 466801 4 0.12 × Comp. Ex.
    18 D 85 7 7 1 0.002 0.006 0.38 0.02 1484 7.6 434475 4 0.12 × Comp. Ex.
    19 D 83 5 6 6 0.002 0.004 0.29 0.04 1497 7.1 411236 4 0.12 × Comp. Ex.
    20 D 83 7 6 4 0.001 0.006 0.33 0.03 1499 6.9 400453 4 0.12 × Comp. Ex.
    21 D 69 17 7 7 0.004 0.009 0.62 0.04 1534 7.8 468633 4 0.14 ⊚ Inventive Ex.
    22 E 62 21 7 10 0.003 0.012 0.47 0.01 1785 12.0 904979 3 0.10 ⊚ Inventive Ex.
    23 F 84 8 6 2 0.004 0.009 0.51 0.01 1328 12.3 595254 4 0.10 ⊚ Inventive Ex.
    24 G 74 13 11 2 0.003 0.013 0.40 0.04 1456 13.0 722246 5 0.10 ○ Inventive Ex.
    25 H 67 4 12 17 0.003 0.010 0.56 0.04 1148 14.3 556223 3 0.12 ⊚ Inventive Ex.
    26 I 54 13 14 19 0.003 0.009 0.60 0.02 1581 8.1 509193 3 0.14 ⊚ Inventive Ex.
    27 J 60 24 9 7 0.001 0.002 0.41 0.04 1337 10.0 488874 4 0.10 Δ Inventive Ex.
    28 K 82 3 10 5 0.001 0.003 0.44 0.03 1352 10.9 541866 3 0.14 ○ Inventive Ex.
    29 L 81 7 11 1 0.005 0.011 0.55 0.04 1480 12.4 706016 4 0.10 ⊚ Inventive Ex.
    30 M 84 1 13 2 0.003 0.012 0.46 0.02 1014 14.8 477880 4 0.08 ○ Inventive Ex.
    31 N 84 4 8 4 0.002 0.006 0.55 0.05 1576 10.2 638167 3 0.12 ⊚ Inventive Ex.
    32 Q 83 0 7 10 0.005 0.012 0.57 0.04 946 8.3 241499 4 0.10 ⊚ Comp. Ex.
    33 P 75 15 8 2 0.003 0.007 0.47 0.02 1494 9.4 542817 4 0.14 × Comp. Ex.
    34 Q 84 2 5 9 0.004 0.012 0.58 0.03 1356 4.8 239679 4 0.10 ⊚ Comp. Ex.
    35 R 52 1 28 19 0.002 0.009 0.47 0.02 1012 28.4 914301 4 0.12 × Comp. Ex.
    36 S 34 1 12 53 0.002 0.006 0.50 0.01 974 14.9 452923 5 0.08 ○ Comp. Ex.
    37 T 64 28 6 2 0.002 0.007 0.44 0.01 1537 6.2 373596 4 0.08 ○ Comp. Ex.
    38 U 63 24 10 3 0.001 0.003 0.45 0.03 1392 11.0 571283 5 0.10 × Comp. Ex.
    39 V 80 3 9 8 0.002 0.004 0.50 0.02 1512 10.7 629088 5 0.12 × Comp. Ex.
    40 W 43 6 9 42 0.002 0.007 0.50 0.01 977 10.3 314542 5 0.14 ⊚ Comp. Ex.
    41 X 49 18 14 19 0.004 0.008 0.44 0.01 1259 17.6 786233 3 0.10 × Comp. Ex.
    42 Y 52 19 9 20 0.000 0.000 - 0.03 1308 9.1 430480 5 0.12 × Comp. Ex.
    43 Z 54 20 8 18 0.062 0.231 0.49 0.01 1504 9.5 554109 4 0.10 × Comp. Ex.
  • The underlines indicate that the values are outside the scope of the present invention. TM: tempered martensite, B: bainite, FM: fresh martensite, γR: retained austenite In the LME evaluation test results, ⊚, ∘, and △ indicate excellent LME resistance, and × indicates poor LME resistance. [Table 3-2]
    Steel sheet microstructure Steel sheet properties LME evaluation test
    Steel sheet sample No. Steel symbol TM + B (%) FM (%) γR (%) Balance (%) Nbsol (mass %) Nbpre (mass %) Formula1 Amount of diffusible hydrogen (mass ppm) TS (MPa) El (%) TS1.5 × El (MPa %) Electrode inclination angle (°) Holding time (s) Evaluation results Remarks
    44 AA 83 2 9 6 0.001 0.005 0.51 0.01 1527 8.2 489297 5 0.12 ⊚ Inventive Ex.
    45 AB 71 20 7 2 0.002 0.006 0.53 0.04 1522 10.8 641277 5 0.10 ⊚ Inventive Ex.
    46 AC 80 1 14 5 0.002 0.009 0.43 0.03 1124 16.2 610470 5 0.12 Δ Inventive Ex.
    47 AD 72 17 11 0 0.002 0.005 0.43 0.02 1178 12.3 497305 5 0.08 Δ Inventive Ex.
    48 AE 80 2 16 2 0.005 0.011 0.49 0.03 1118 18.8 702782 4 0.10 ○ Inventive Ex.
    49 AF 63 24 10 3 0.003 0.008 0.49 0.03 1308 9.4 444672 4 0.12 ⊚ Inventive Ex.
    50 AG 84 1 8 7 0.001 0.003 0.47 0.01 1292 9.1 422606 3 0.14 Δ Inventive Ex.
    51 AH 79 4 9 8 0.003 0.007 0.48 0.01 1321 9.6 460920 3 0.12 ⊚ Inventive Ex.
    52 AI 80 3 17 0 0.003 0.007 0.46 0.04 1116 20.5 764276 4 0.12 Δ Inventive Ex.
    53 AJ 80 1 11 8 0.002 0.006 0.44 0.03 1261 11.7 523912 5 0.08 ○ Inventive Ex.
    54 AK 81 4 14 1 0.002 0.007 0.42 0.04 1152 13.3 520032 3 0.08 ○ Inventive Ex.
    55 AL 79 4 15 2 0.003 0.008 0.42 0.02 1090 18.8 676546 3 0.12 Δ Inventive Ex.
    56 AM 77 2 13 8 0.001 0.005 0.43 0.03 1300 13.6 637461 3 0.12 Δ Inventive Ex.
    57 AN 79 6 12 3 0.002 0.006 0.49 0.03 1257 13.1 583814 4 0.08 ⊚ Inventive Ex.
    58 AO 71 16 11 2 0.003 0.009 0.41 0.03 1130 12.5 474819 3 0.10 ○ Inventive Ex.
    59 AP 47 24 10 19 0.003 0.009 0.47 0.04 1265 11.8 530906 4 0.10 ○ Inventive Ex.
    60 AQ 52 20 10 18 0.003 0.008 0.41 0.02 1112 12.4 459810 4 0.14 ⊚ Inventive Ex.
    61 AR 83 1 12 4 0.004 0.009 0.46 0.02 1087 14.0 501733 3 0.12 ⊚ Inventive Ex.
    62 AS 80 8 6 6 0.002 0.010 0.46 0.02 1425 7.4 398065 5 0.12 ○ Inventive Ex.
    63 AT 64 24 8 4 0.003 0.008 0.55 0.02 1398 8.2 428622 3 0.08 ⊚ Inventive Ex.
    64 AU 57 25 7 11 0.004 0.010 0.88 0.03 1636 9.5 628635 5 0.10 ○ Inventive Ex.
    65 AV 59 24 12 5 0.005 0.006 0.42 0.05 1374 14.4 733403 3 0.08 ⊚ Inventive Ex.
    66 AW 68 17 9 6 0.005 0.009 0.54 0.02 1399 13.1 685485 5 0.10 ○ Inventive Ex.
    67 AX 61 22 5 12 0.092 0.101 0.41 0.05 1639 6.8 451209 4 0.12 Δ Inventive Ex.
    68 AY 66 18 9 7 0.005 0.010 0.63 0.04 1360 11.4 571759 4 0.08 ○ Inventive Ex.
    69 AZ 84 4 9 3 0.003 0.010 0.81 0.03 1502 10.9 634500 3 0.10 ○ Inventive Ex.
    70 E 63 22 6 9 0.003 0.013 0.84 0.46 1798 10.6 808147 5 0.10 Δ Inventive Ex.
  • The underlines indicate that the values are outside the scope of the present invention. TM: tempered martensite, B: bainite, FM: fresh martensite, yR: retained austenite In the LME evaluation test results, ⊚, ∘, and △ indicate excellent LME resistance, and × indicates poor LME resistance.
  • The area fractions of fresh martensite, tempered martensite, and bainite were determined by the following method. A thickness-direction cross-section (L cross-section) of each steel sheet, which was parallel to the rolling direction of the sheet, was polished and then etched with Nital. The cross-section was observed by a scanning electron microscope (SEM) at 2000-fold magnification in 10 fields of view at a 1/4 thickness position (a position at a distance of 1/4 of the sheet thickness from the steel sheet surface). Using a microstructural image obtained, the area fractions of constituent microstructures (the total fraction of tempered martensite and bainite, and the fraction of fresh martensite) were calculated. In the microstructural image, fresh martensite was determined by a light gray microstructural region, and tempered martensite and bainite were determined by a dark gray region in which carbides were precipitated.
  • The area fraction of retained austenite was determined by the following method. Each steel sheet was polished from the 1/4 thickness position to a plane by a thickness of 0.1 mm, and then further chemically polished to a plane by a thickness of 0.1 mm. The plane (surface) was subjected to X-ray diffraction analysis using CoKα rays to measure the integrated intensity ratios of the diffraction peaks of the {200}, {220} and {311} planes of fcc iron, and the {200}, {211} and {220} planes of bcc iron. The nine integrated intensity ratios obtained were averaged to determine the area fraction of retained austenite.
  • The amount of hydrogen in steel was determined by the following method. Test specimens of about 5 × 30 mm were cut out from a hot-rolled steel sheet, a cold-rolled steel sheet, and a galvanized steel sheet. For the test specimen of galvanized steel sheet, the coating on the surface of the test specimen was removed in advance using a router (precision grinder). Each test specimen was placed in a quartz tube, and the internal atmosphere of the tube was replaced with Ar, and then heated at 200°C/hr. Hydrogen generated until the temperature reached 400°C was subjected to gas chromatography to measure the amount of released hydrogen by a temperature rise analysis method. The cumulative value of the amount of hydrogen detected in the temperature range from room temperature (25°C) to less than 250°C was taken as the amount of diffusible hydrogen.
  • The total amount of Nb contained in the steel (Nbtotal), the amount of dissolved Nb in the steel sheet (Nbsol), and the amount of Nb in Nb precipitates having a particle size of less than 20 nm (Nbpre) were measured in the following manner. The total amount of Nb contained in the steel (Nb) was measured by wet chemical analysis. The amount of Nb in Nb precipitates having a particle size of 20 nm or less was measured by the following method. Precipitates in the steel material were captured as residues, and the amount of Nb in all residues was determined. Thereafter, the amount of Nb present in residues having a particle size of 20 nm or more was determined. The amount of Nb in Nb precipitates having a particle size of 20 nm or less was determined as the difference between the amount of Nb in all residues and the amount of Nb present in the residues having a particle size of 20 nm or more. Specific procedures are as follows. A plurality of test specimens of about 20 × 50 mm were cut out from a hot-rolled steel sheet, a cold-rolled steel sheet, or a galvanized steel sheet. For a test specimen of a galvanized steel sheet, the coating on the surface of the test specimen was removed in advance using a router (precision grinder). The surface of each test specimen was polished about 50 µm by preliminary electrolytic polishing to obtain a fresh surface. The resulting test specimen was subjected to electrolysis using 10 vol % acetylacetone - 1 mass % tetramethylammonium chloride - methanol as an electrolytic solution for extracting precipitates. The resulting electrolytic solution after electrolysis was passed through a filter having a pore size of 0.2 µm to capture residues. The residues were then decomposed with an acid, and the Nb concentration was quantified in mass % unit using ICP emission spectrometry. The value obtained was used as the amount of Nb in all residues. Subsequently, a remaining test specimen was subjected to electrolysis using 10 vol % acetylacetone - 1 mass % tetramethylammonium chloride - methanol. Thereafter, the residual portion of the metal sample, remaining after the electrolysis, was immersed in separately prepared methanol, and residues, etc. attached to the residual portion of the metal sample were collected in a container using ultrasonic vibration. Thereafter, the electrolytic solution after electrolysis and the methanol containing the residues that had been attached to the residual portion of the metal sample were passed through an alumina filter having a pore size of 20 nm to capture residues. The residues were immersed in hexametaphosphoric acid and dispersed in the hexametaphosphoric acid using ultrasonic vibration. The residues, dispersed in hexametaphosphoric acid, were supplemented using a new alumina filter having a pore size of 20 nm. The residues captured on the alumina filter had a particle size of 20 nm or more. The residues were decomposed with an acid, and the Nb concentration was quantified in mass % unit using ICP emission spectrometry. The value obtained was used as the amount of Nb present in the residues having a particle size of 20 nm or more. The difference between the amount of Nb in all residues and the amount of Nb present in the residues having a particle size of 20 nm or more was taken as the amount of Nb in Nb precipitates having a particle size of 20 nm or less (Nbpre). Further, the amount of dissolved Nb in the steel sheet (Nbsol) was determined as the difference between the total amount of Nb contained in the steel (Nb) and the amount of Nb in all residues determined by the above method.
  • While the above-described Nb precipitates are composed mainly of NbC, an NbN precipitate or other Nb precipitates may be contained in the steel sheet.
  • A tensile test was performed according to JIS Z 2241 (2011) using a JIS No. 5 test specimen which had been taken from a steel sheet such that the tensile direction was perpendicular to the rolling direction of the steel sheet, and the TS (tensile strength) and EL (total elongation) of the test specimen were measured. In the Examples, steel sheets having a TS of 980 MPa or more were evaluated as acceptable, judging that the intended strength was achieved. Steel sheets having a "TS1.5 × El" of 390,000 or more were evaluated as acceptable, judging that the intended formability was achieved.
  • The LME resistance was evaluated in the following manner. A sample, having the dimensions of: 100 mm in a direction perpendicular to the rolling direction, and 30 mm in the rolling direction, was taken from a steel sheet. The sample and a 980GA sample having the same size were stacked to prepare an evaluation sample.
  • Resistance spot welding was performed on the evaluation sample under the conditions of an electrode inclination angle of θ and a welding pressure of 3.5 kN. The electrode inclination angle in spot welding is herein defined as the angle θ between a line passing through the major axis of a nugget and a line parallel to the surface of a steel sheet in a cross-section of a spot-welded member. The welding current pattern was controlled so that the diameter of a nugget formed would fall within the range of 3.5√t to 5.5√t, t being the thickness (1.2 mm) of one steel sheet. In the resistance spot welding, Dr6-type CuCr electrodes were used, and the clearance between the evaluation sample and an electrode was 2.0 mm. The electrode inclination angles and holding times in the LME resistance evaluation tests are shown in Tables 3-1 and 3-2.
  • In each evaluation test, 5 evaluation samples were spot-welded to produce welded members, and a cross-section of each spot-welded member was observed using an optical microscope (100-fold magnification) to evaluate the LME resistance. When the number of members with cracking was two or less, and the average cracking depth was less than 100 µm (including no occurrence of cracking), the steel sheet was rated "⊚", when the number of members with cracking was two or less, and the average cracking depth was not less than 100 µm and not more than 300 µm, the steel sheet was rated "∘", when the number of members with cracking was two or less, and the average cracking depth was 300 µm or more, the steel sheet was rated "△", and when the number of members with cracking was three or more, the steel sheet was rated "×". In the Examples, the steel sheets rated ⊚, ∘, and △ were judged to have excellent LME resistance and referred to as Inventive Examples, while the steel sheets rated × were judged to have poor LME resistance and referred to as Comparative Examples.
  • The high-strength steel sheets of the Inventive Examples all have a high TS and a good TS × El balance, and are excellent in LME resistance, whereas the steel sheets of the Comparative Examples are poor in at least one of TS, TS × El balance, and LME resistance.

Claims (8)

  1. A high-strength steel sheet having a chemical composition comprising, in mass %:
    C: not less than 0.030% and not more than 0.500%,
    Si: more than 0.01% and not more than 2.50%,
    Mn: not less than 0.10% and not more than 5.00%,
    P: 0.100% or less,
    S: 0.0200% or less,
    Al: 0.100% or less,
    N: 0.0100% or less,
    O: 0.0100% or less,
    Ti: not less than 0.010% and not more than 0.200%, and
    Nb: not less than 0.005% and not more than 0.500%,
    with the balance being Fe and incidental impurities,
    wherein in the microstructure of the steel sheet at a 1/4 thickness position,
    the total area fraction of tempered martensite and bainite is not less than 40% and not more than 85%,
    the area fraction of fresh martensite is not less than 0% and not more than 25%,
    the area fraction of retained austenite is not less than 5% and not more than 20%, and
    the balance includes at least one of ferrite and pearlite at an area fraction of not less than 0% and not more than 20%,
    wherein the relationship between the amount of dissolved Nb (Nbsol), the amount of Nb in Nb precipitates having a particle size of less than 20 nm (Nbpre), and the total amount of Nb contained in the steel sheet (Nb) satisfies the following Formula 1, and
    wherein the amount of diffusible hydrogen in the steel is 0.50 mass ppm or less: Nb sol / Nb + Nb pre / Nb ≥ 0.40 where Nbsol represents the amount (mass %) of dissolved Nb, and Nbpre represents the amount (mass %) of Nb in Nb precipitates having a particle size of less than 20 nm.
  2. The high-strength steel sheet according to claim 1, wherein the chemical composition further comprises, in mass %, one or two or more selected from the following:
    V: 0.500% or less,
    Ta: 0.10% or less,
    W: 0.10% or less,
    B: 0.0100% or less,
    Cr: 1.00% or less,
    Mo: 1.00% or less,
    Co: 1.00% or less,
    Ni: 1.00% or less,
    Cu: 1.00% or less,
    Sn: 0.200% or less,
    Sb: 0.200% or less,
    Ca: 0.0100% or less,
    Mg: 0.0100% or less,
    REM: 0.0100% or less,
    Zr: 0.100% or less,
    Te: 0.100% or less,
    Hf: 0.10% or less, and
    Bi: 0.200% or less.
  3. The high-strength steel sheet according to claim 1 or 2, having a coating layer on the surface of the steel sheet.
  4. The high-strength steel sheet according to claim 3, wherein the coating layer is an alloyed coating layer.
  5. A method for manufacturing the high-strength steel sheet according to claim 1 or 2, comprising:
    a slab heating step of heating a steel material having the chemical composition at a temperature equal to or higher than Tsol°C represented by the following Formula 2 for 1.0 hours or more;
    a hot rolling step;
    a coiling step at a coiling temperature TCT°C of 650°C or lower; and
    an annealing step of heating a steel sheet to a soaking temperature TAT°C, holding the steel sheet at the soaking temperature TAT°C of not less than 750°C and not more than 950°C, and then cooling the steel sheet,
    wherein QHR defined by Formula 3 in terms of a finish rolling start temperature TFET°C, a finish rolling delivery temperature TFDT°C, and an effective time tHR from the start of finish rolling to the end of finish rolling in the hot rolling step,
    QCT defined by Formula 4 in terms of a residence time tCT from TFDT°C to 650°C and the coiling temperature TCT°C in the coiling step, and
    QAT defined by Formula 5 in terms of the soaking temperature TAT°C, a heating time t from 650°C to the soaking temperature TAT°C, and a holding time tAT at the soaking temperature TAT°C in the annealing step, satisfy Formula 6: T sol = 7900 / 3.42 − log Nb % C % − 273 where [Nb%] and [C%] respectively represent the Nb content (mass %) and the C content (mass %) of the steel; Q HR = 0.5 T FET + T FDT × log 10 t HR ; Q CT = 0.5 T CT + 650 × log 10 t CT ; QAT = 0.5(650 + TAT) × log10(t) + TAT × log10(tAT); and Q HR + Q CT + Q AT ≤ 6000 .
  6. The method for manufacturing a high-strength steel sheet according to claim 5, wherein the value of QAT is 4700 or less.
  7. The method for manufacturing a high-strength steel sheet according to claim 5 or 6, wherein the steel sheet after the annealing step is subjected to a coating treatment.
  8. The method for manufacturing a high-strength steel sheet according to claim 7, wherein the coating treatment is an alloying coating treatment.
EP24810653.6A 2023-05-19 2024-02-27 High-strength steel sheet and method for producing same Pending EP4685260A1 (en)

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WO2020184154A1 (en) 2019-03-11 2020-09-17 Ｊｆｅスチール株式会社 High-strength steel sheet and method for producing same
WO2020225936A1 (en) 2019-05-09 2020-11-12 日本製鉄株式会社 Steel sheet and method for manufacturing same

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JPS6456853A (en) * 1987-08-28 1989-03-03 Nippon Kokan Kk Square tube excellent in molten zinc cracking resistance
JP5283402B2 (en) * 2008-03-07 2013-09-04 日新製鋼株式会社 Zn-Al-Mg plated steel sheet with excellent resistance to molten metal embrittlement cracking
WO2018234839A1 (en) * 2017-06-20 2018-12-27 Arcelormittal Zinc coated steel sheet with high resistance spot weldability
MX2023011867A (en) * 2021-04-09 2023-10-23 Jfe Steel Corp High-strength cold-rolled steel sheet and method for manufacturing same.
MX2023009752A (en) * 2021-04-27 2023-08-30 Nippon Steel Corp Steel sheet and plated steel sheet.
EP4361303A4 (en) * 2021-07-09 2024-11-06 JFE Steel Corporation HIGH-STRENGTH STEEL SHEET, HIGH-STRENGTH PLATED STEEL SHEET, MANUFACTURING METHOD THEREOF, AND MEMBER

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Publication number Priority date Publication date Assignee Title
WO2016159169A1 (en) 2015-03-30 2016-10-06 新日鐵住金株式会社 Method for spot welding of plated steel sheet
WO2020184154A1 (en) 2019-03-11 2020-09-17 Ｊｆｅスチール株式会社 High-strength steel sheet and method for producing same
WO2020225936A1 (en) 2019-05-09 2020-11-12 日本製鉄株式会社 Steel sheet and method for manufacturing same

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See also references of WO2024241645A1

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