EP3604587A1 - Warmgewalztes stahlblech, stahlblechteil und herstellungsverfahren dafür - Google Patents

Warmgewalztes stahlblech, stahlblechteil und herstellungsverfahren dafür Download PDF

Info

Publication number
EP3604587A1
EP3604587A1 EP17904326.0A EP17904326A EP3604587A1 EP 3604587 A1 EP3604587 A1 EP 3604587A1 EP 17904326 A EP17904326 A EP 17904326A EP 3604587 A1 EP3604587 A1 EP 3604587A1
Authority
EP
European Patent Office
Prior art keywords
steel sheet
less
rolling
rem
content
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.)
Withdrawn
Application number
EP17904326.0A
Other languages
English (en)
French (fr)
Other versions
EP3604587A4 (de
Inventor
Tatsuo Yokoi
Nobuo Yoshikawa
Shigeru Yonemura
Kazuya Ootsuka
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.)
Nippon Steel Corp
Original Assignee
Nippon 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 Nippon Steel Corp filed Critical Nippon Steel Corp
Publication of EP3604587A1 publication Critical patent/EP3604587A1/de
Publication of EP3604587A4 publication Critical patent/EP3604587A4/de
Withdrawn legal-status Critical Current

Links

Images

Classifications

    • 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
    • C21D1/02—Hardening articles or materials formed by forging or rolling, with no further heating beyond that required for the formation
    • 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21C—PROCESSING OF PIG-IRON, e.g. REFINING, MANUFACTURE OF WROUGHT-IRON OR STEEL; TREATMENT IN MOLTEN STATE OF FERROUS ALLOYS
    • C21C7/00—Treating molten ferrous alloys, e.g. steel, not covered by groups C21C1/00 - C21C5/00
    • C21C7/04—Removing impurities by adding a treating agent
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21C—PROCESSING OF PIG-IRON, e.g. REFINING, MANUFACTURE OF WROUGHT-IRON OR STEEL; TREATMENT IN MOLTEN STATE OF FERROUS ALLOYS
    • C21C7/00—Treating molten ferrous alloys, e.g. steel, not covered by groups C21C1/00 - C21C5/00
    • C21C7/04—Removing impurities by adding a treating agent
    • C21C7/06—Deoxidising, e.g. killing
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/19—Hardening; Quenching with or without subsequent tempering by interrupted quenching
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • 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
    • 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
    • C21D8/0215—Rapid solidification; Thin strip casting
    • 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
    • C21D8/0221—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
    • 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
    • C21D8/04—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 to produce plates or strips for drawing, e.g. for deep-drawing
    • C21D8/041—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 to produce plates or strips for drawing, e.g. for deep-drawing involving a particular fabrication or treatment of ingot or slab
    • C21D8/0415—Rapid solidification; Thin strip casting
    • 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
    • C21D8/04—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 to produce plates or strips for drawing, e.g. for deep-drawing
    • C21D8/0421—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 to produce plates or strips for drawing, e.g. for deep-drawing characterised by the working steps
    • C21D8/0426—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties 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
    • C21D8/04—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 to produce plates or strips for drawing, e.g. for deep-drawing
    • C21D8/0447—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 to produce plates or strips for drawing, e.g. for deep-drawing characterised by the heat treatment
    • C21D8/0463—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 to produce plates or strips for drawing, e.g. for deep-drawing 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
    • 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
    • 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
    • C21D9/48—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals deep-drawing sheets
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—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
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/08—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
    • C22C38/12—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
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing 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
    • C22C38/16—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
    • C22C38/32—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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
    • 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
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/009—Pearlite

Definitions

  • the present invention relates to a hot rolled steel sheet and a steel forged part, and a production method therefor.
  • High strength and high press workability are required for steel sheets used in body structures of automobiles in view of safety improvement and weight reduction.
  • a high-strength steel sheet that has excellent hole expandability a high burring property
  • ⁇ value a steel sheet made, as a primary phase, of ferrite that is precipitation strengthened with fine precipitates such as Ti and Nb and a production method therefor have been reported.
  • Patent Document 1 discloses a hot-rolled steel sheet that is strong and has excellent stretch flangeability.
  • Patent Document 2 discloses a high-formability high-tensile-strength hot-rolled steel sheet that has excellent material uniformity.
  • Patent Document 3 discloses a high tensile strength hot-rolled steel sheet that has excellent elongation and stretch flangeability.
  • the sheet metal forging is a way of press working that includes mixed working elements including forging-specific working elements, in addition to conventional working elements for press working steel sheets.
  • a steel sheet is deformed into a shaped part with the steel sheet retaining an original sheet thickness or being thinned (reduced in thickness) by the conventional press working, while the sheet thickness is increased in a forged portion by a partially applied compressive force.
  • efficient deformation can be achieved such that a sheet thickness of the steel sheet intended for a functionally necessary portion can be attained, and strength of the part can be secured.
  • Patent Documents 1 to 3 fail to describe a working including mixed working elements typified by the sheet metal forging. Further, coiling conditions for producing a hot-rolled steel sheet described in Patent Document 1 are highly strict and are not practical. Still further, the hot-rolled steel sheets described in Patent Documents 2 and 3 contain Mo, which is an expensive alloying element, in an amount of 0.07% or more, and thus have a problem of high production costs.
  • press cracking appears at a point where sheet thickness necking (a reduced sheet thickness of the steel sheet) occurs. It has also been found that even in a working that is not associated with sheet thickness necking, such as sheet metal forging, cracks may be generated in the material, which may end in rupture and products may not be obtained in some cases.
  • An object of the present invention which has been made to solve the above problem, is to provide a hot rolled steel sheet with excellent sheet forgeability, which maintains basic features as a high burring steel and also makes it possible to improve cracking limit of a forged portion by a partially applied compressive force.
  • the present invention has been made to solve the above problem, and the gist thereof a hot rolled steel sheet and a steel forged part, and a production method therefor, as described below.
  • a hot rolled steel sheet with excellent sheet forgeability maintaining a good hole expandability which is a basic feature as a high burring steel, can be provided.
  • Figure 1 shows schematic drawings illustrating a simple shear test.
  • Figure 1(a) illustrates a specimen for a simple shear test.
  • Figure 1(b) illustrates a specimen after a simple shear test.
  • the present inventors conducted intensive studies in order to solve the above problem and obtained the following findings.
  • the sheet metal forging includes a strain range exceeding a rupture strain in a conventional tensile test (high strain range). Since the sheet metal forging is a composite working, it cannot be evaluated simply based on tensile test and shear test data. Accordingly, the present inventors established a new way of evaluation by introducing an "equivalent plastic strain" as an indicator.
  • the equivalent plastic strain can be used as an indicator to mixedly evaluate a tensile stress and a tensile strain at the time of rupture when a tensile test is conducted and a shearing stress and a shearing strain at the time of rupture when a shear test is conducted.
  • the equivalent plastic strain is converted using a relation between a shearing stress ⁇ s and a shear plastic strain ⁇ sp in a simple shear test into a relation between a tensile stress ⁇ and a tensile strain ⁇ in a uniaxial tensile test, which is different in deformation mode.
  • a constant, conversion factor ( ⁇ ) can be used to make a conversion as in the formula below.
  • the conversion factor ( ⁇ ) is calculated according to a method described later, and then an equivalent plastic strain is derived.
  • the sheet metal forging includes deformation in a high strain range. Accordingly, when test is performed at one time in a commonly used shear test device, cracks may propagate in a specimen from a portion where the specimen is held. As a result, a test of deformation may not often be completed up to the high strain range. Therefore, there is a need for a method for reproducing a working, such as sheet metal forging, in which thinning (thickness reduction and necking) of steel sheet does not occur.
  • the present inventors have then chosen to divide a shear test into multiple stages, machine an initiation point of a crack in a specimen generated in a portion where the specimen is held in order to prevent the crack from propagating in the specimen after the shear test of each stage, and evaluate a test result obtained by serially connecting the shear test results.
  • Employing the test method it is possible to obtain the shear test results up to the high strain range and to determine a relation between a shearing stress and a shearing strain up to the high strain range.
  • a conventional tensile test method can be applied to the tensile stress and the tensile strain.
  • a JIS No. 5 specimen based on JIS Z 2241 (2011) can be used.
  • TiN precipitated TiN
  • ferrite that is precipitation strengthened with fine precipitates such as Ti and Nb (precipitation strengthened ferrite)
  • Ti and Nb precipitation strengthened ferrite
  • coarse TiN may be precipitated (hereafter, such precipitated TiN is also referred to simply as "TiN") unless any special production method is used. This is because TiN is a highly thermally-stable compound from a thermodynamic point of view and may preferentially be crystalized or precipitated with respect to other compounds when being cast during a steel sheet production process, when heated in hot rolling, or under a high temperature condition as with the beginning of rough rolling.
  • TiN is hard enough to be used for cutting tools, machine parts, press tooling for plastic molding, sporting goods, coating applications for ornaments, and the like. It is known that the hardness of TiN is on the order of Hv2000 to 2300, which demonstrates that TiN is a very hard precipitate. Accordingly, under deformation in a high strain range as in the case of sheet metal forging, a void is likely to be generated at an interface between a parent phase structure and TiN due to a difference in deformability with the parent phase structure.
  • a void may be generated at an interface between the two phases. Thereafter, as strain associated with the sheet metal forging increases, the void may grow and coalesce with an adjacent void to become a crack, ending in rupture.
  • the present inventors have then found that the crack generation can be inhibited if the void generation can be prevented and if the void can be inhibited from coalescing with an adjacent void even when the void grows. At this time, however, it is also important that intrinsic functionality as a high burring steel is left unimpaired.
  • the present inventors have found the followings from the findings.
  • a void may be generated at a grain boundary of TiN, which is a hard precipitate, and thus limiting an average diameter of TiN can lead to a reduction in void generation.
  • a void may be generated at a grain boundary of TiN, and thus spacing the TiN precipitates apart from each other can make it difficult for voids to coalesce with each other even when the voids grow.
  • the void generation can be reduced by reducing a difference in hardness between a hard microstructure and a soft microstructure as much as possible.
  • the effective cumulative strain is an indicator that takes into consideration grain recovery, recrystallization, and grain growth according to temperature during rolling and rolling reduction of a steel sheet by rolling. Accordingly, a constitutive equation that represents static recovery phenomena in a time lapse after rolling is used for determining the effective cumulative strain.
  • the static recovery of grains in a time lapse after rolling is taken into consideration because energy accumulated as strain in rolled grains may be released in the static recovery due to vanishment of thermal dislocations of grains. Further, the vanishment of thermal dislocations may be affected by rolling temperature and lapsed time after rolling. Accordingly, taking the static recovery into consideration, the present inventors introduced an indicator described, as parameters, by the temperature during rolling, the rolling reduction of a steel sheet by rolling (logarithmic strain), and the lapsed time after rolling, and defined it as "effective cumulative strain”.
  • complex oxide including Ca, Mg, or REM and including any one or more of Ti and Al (hereafter, also referred to simply as "complex oxide") in steel during deoxidation process.
  • complex oxide including Ca, Mg, or REM and including any one or more of Ti and Al (hereafter, also referred to simply as "complex oxide") in steel during deoxidation process.
  • the present inventors have found that this can only be achieved by sequentially adding strong deoxidizing elements in the deoxidation process.
  • strong deoxidizing elements are added to molten steel in which a weak deoxidizing element oxide is present.
  • the weak deoxidizing element oxide may be reduced by the strong deoxidizing elements, and thus oxygen is released at a slow supply rate and low supersaturation.
  • an oxide generated from the added strong deoxidizing elements may be fine.
  • the deoxidation method maximizes these effects by adding deoxidizing elements stepwise: starting from Si, which is a weak deoxidizing element, then Ti, Al, sequentially, and then Ca, Mg, or REM, which is a strong deoxidizing element.
  • C carbon
  • elements such as Nb and Ti
  • a content of C is less than 0.020%, a sufficient effect of the action cannot be obtained.
  • more than 0.070% of the content of C may cause an increase in iron-carbides, which may be an initiation point of cracking during hole expansion working, and thus the hole expansion value is degraded.
  • the content of C is 0.020 to 0.070%.
  • the content of C is preferably 0.025% or more, and more preferably 0.030% or more.
  • the content of C is preferably 0.060% or less, and more preferably 0.050% or less.
  • Si silicon
  • Si has a deoxidation effect, and an effect of inhibiting precipitation of iron-carbides such as cementite in a material microstructure and contributing to improvement in ductility and hole expandability.
  • ferritic transformation may be likely to occur in a high temperature range, and thus carbides including Ti may be likely to precipitate.
  • the precipitation of carbide in a high temperature range is likely to cause variation in the amount of precipitation, resulting in material variation in strength, hole expandability, and the like. Accordingly, a content of Si is 0.05 to 1.70%.
  • the content of Si is preferably 0.06% or more, and more preferably 0.08% or more.
  • the content of Si is preferably 1.50% or less, and in view of improving chemical treatability and a corrosion resistance after painting, the content of Si is more preferably 1.00% or less.
  • Mn manganese
  • Mn manganese
  • a content of Mn is 0.60 to 2.50%.
  • the content of Mn is preferably 1.00% or more, and more preferably 1.50% or more.
  • the content of Mn is preferably 2.00% or less, and more preferably 1.80% or less.
  • Al (aluminum) has a deoxidation effect and an effect of generating ferrite, as with Si. On the other hand, an excessive content may lead to embrittlement and be likely to cause clogging of a tundish nozzle during casting. Accordingly, a content of Al is 0.005 to 0.020%.
  • the content of Al is preferably 0.007% or more, and more preferably 0.010% or more.
  • the content of Al is preferably 0.018% or less, and more preferably 0.015% or less.
  • N more than 0.0030% to 0.0060% or less
  • N nitrogen
  • nitrogen nitrogen
  • the coarsened TiN may serve as an initiation point of a void and promote fracture.
  • the content of N is more than 0.0030% to 0.0060% or less.
  • the content of N is preferably 0.0.0031% or more, and more preferably 0.0035% or more.
  • the content of N is preferably 0.0055% or less, and more preferably 0.0050% or less.
  • P phosphorus
  • a content of P is preferably as small as possible. Accordingly, the content of P is limited to 0.050% or less.
  • the content of P is preferably 0.030% or less or 0.020% or less. It is not particularly necessary to define a lower limit, and the lower limit is 0%. However, an excessive reduction in the content of P leads to an increase in costs during smelting, and thus the lower limit may be 0.001%.
  • S sulfur
  • S is also an impurity contained in molten pig iron, and since S may degrade local ductility and weldability due to formation of MnS, a content of S is preferably as small as possible. Accordingly, the content of S is limited to 0.005% or less. To improve ductility and weldability, the content of S may be 0.003% or less or 0.002% or less. It is not particularly necessary to define a lower limit, and the lower limit is 0%. However, an excessive reduction in the content of S leads to an increase in costs during smelting, and thus the lower limit may be 0.0005%.
  • Ti has an effect of improving low temperature toughness because carbo-nitride or dissolved Ti may cause a delay in grain growth during hot rolling and thus refine grain diameter in a hot rolled sheet. Further, Ti may be finely dispersed in ferrite as TiC, so that it contributes to strengthening of the steel sheet through precipitation strengthening. However, an excessive content may cause saturation of the effect and make TiN, which is a hard precipitate, to likely to precipitate. Accordingly, a content of Ti is 0.015 to 0.170%. The content of Ti is preferably 0.030% or more, 0.045% or more or 0.060% or more, and more preferably 0.070% or more, 0.080% or more, 0.090% or more, or 0.100% or more. In addition, the content of Ti is preferably 0.160% or less, 0.150% or less, 0.140% or less, 0.130% or less, or 0.120% or less.
  • O oxygen
  • Oxgen is an element necessary for dispersing a number of fine oxides during molten steel deoxidation, and 0.0010% or more of O may be contained in a final product steel in order to produce such an effect.
  • more than 0.0100% of O content may cause not only saturation in the effect but also an increase in size and density of oxides as an impurity.
  • the oxide may serve as an initiation point of a void and promote fracture.
  • a content of O is 0.0010 to 0.0100%.
  • the content of O is preferably 0.0030% or more, and more preferably 0.0050% or more.
  • the content of O is preferably 0.0090% or less, and more preferably 0.0080% or less.
  • Nb (niobium) has an effect of improving low temperature toughness because carbo-nitride or dissolved Nb may cause a delay in grain growth during hot rolling and thus refine grain diameter in a hot rolled sheet.
  • Nb may be present as NbC, so that it contributes to strengthening of the steel sheet through precipitation strengthening. Accordingly, Nb may be contained as necessary. However, an excessive content may cause saturation of the effect, leading to a decrease in economy. Accordingly, a content of Nb is 0.100% or less.
  • the content of Nb may be 0.080% or less, 0.060% or less, or 0.050% or less, as necessary.
  • a lower limit of Nb is 0%. However, the lower limit may be 0.001% or 0.010% in order to produce the effect sufficiently.
  • V vanadium
  • V vanadium
  • the content of V may be 0.200% or less, 0.100% or less, or 0.060% or less, as necessary.
  • a lower limit of Nb is 0%. However, the lower limit may be 0.001% or 0.010% in order to produce the effect sufficiently.
  • Cu copper is an element that has an effect of improving strength of a steel sheet by precipitation strengthening or solid solution strengthening. Accordingly, Cu may be contained as necessary. However, an excessive content may cause saturation of the effect, leading to a decrease in economy. Accordingly, a content of Cu is 2.00% or less. Further, a large amount of Cu content may cause a blemish due to a scale on a surface of the steel sheet. Accordingly, the content of Cu may be 1.20% or less, 0.80% or less, 0.50% or less, or 0.25% or less. A lower limit of Cu is 0%. However, the lower limit of Cu may be 0.01% in order to produce the effect sufficiently.
  • Ni nickel
  • Ni is an element that has an effect of improving strength of a steel sheet by solid solution strengthening. Accordingly, Ni may be contained as necessary. However, an excessive content may cause saturation of the effect, leading to a decrease in economy. Accordingly, a content of Ni is 2.00% or less. Further, a large amount of Ni content may cause degradation of ductility. Accordingly, the content of Ni may be 0.60% or less, 0.35% or less, or 0.20% or less. A lower limit of Ni is 0%. However, the lower limit of Ni may be 0.01% in order to produce the effect sufficiently.
  • Cr chromium
  • Cr is an element that has an effect of improving strength of a steel sheet by solid solution strengthening. Accordingly, Cr may be contained as necessary. However, an excessive content may cause saturation of the effect, leading to a decrease in economy. Accordingly, a content of Cr is 2.00% or less. To improve economy, an upper limit of Cr may be 1.00%, 0.60%, or 0.30%. A lower limit of Cr is 0%. However, the lower limit of Cr may be 0.01% in order to produce the effect sufficiently.
  • Mo mobdenum
  • Mo is an element that has an effect of improving strength of a steel sheet by precipitation strengthening or solid solution strengthening. Accordingly, Mo may be contained as necessary. However, an excessive content may cause saturation of the effect, leading to a decrease in economy. Accordingly, a content of Mo is 1.00% or less. To improve economy, an upper limit of Mo may be 0.60%, 0.30%, or 0.10%. A lower limit of Mo is 0%. However, the lower limit of Mo may be 0.005% or 0.01% in order to produce the effect sufficiently.
  • B (boron) segregates at a grain boundary, and may increase grain boundary strength to improve low temperature toughness. Accordingly, B may be contained as necessary. However, an excessive content may cause saturation of the effect, leading to a decrease in economy. Accordingly, a content of B is 0.0100% or less. Further, B is a strong quench-hardening element, and a large amount of B content may prevent ferritic transformation from sufficiently progressing during cooling and sufficient retained austenite may not be obtained. Accordingly, a content of B may be 0.0050% or less, 0.0020% or less, or 0.0015%. A lower limit of B is 0%. However, the lower limit of B may be 0.0001% or 0.0002% in order to produce the effect sufficiently.
  • Ca, Mg, and REM are strong deoxidizing elements. As described above, by adding these elements to molten steel subsequently to Ti and Al, followed by sequential deoxidation, oxides are made fine. As a result, TiN precipitated with these elements being cores is also made fine, and a high-strength hot-rolled steel sheet excellent for sheet metal forging can be obtained.
  • one or more selected from Ca, Mg, and REM are to be contained so that a total of the contents satisfies the above formula (i).
  • the content of Ca or Mg is more than 0.0100% or the content of REM is more than 0.1000%, not only the effect saturates, but also the size and the density of oxides containing Al or Ca increase.
  • the oxide containing Al or Ca may serve as an initiation point of a void and promote fracture.
  • contents of Ca and Mg are 0.0100% or less, a content of REM is 0.1000% or less.
  • a total of the contents is preferably 0.0010% or more, 0.0015% or more, or 0.0020% or more.
  • the content of each of Ca and Mg is preferably 0.0090% or less, 0.0080% or less, or 0.0070% or less.
  • the content of REM is preferably 0.0900% or less, 0.0800% or less, or 0.0700% or less.
  • REM refers to a total of 17 elements of Sc, Y and lanthanoid, and the content of REM means a total content of these elements. It is to be noted that lanthanoid is industrially added in the form of a mischmetal.
  • Zr, Co, Zn, and W are each 1.000% or less, the effect of the present invention is unimpaired even if contained.
  • An upper limit of each of them may be 0.300% or 0.10%.
  • a total content of Zr, Co, Zn, and W is preferably 1.000% or less, or 0.100%. These elements may not necessarily be contained, and a lower limit is 0%, although the lower limit may be 0.0001% as necessary.
  • the effect of the present invention is unimpaired if a small amount of Sn (tin) is contained.
  • the content of more than 0.05% may be a cause of a flaw during hot rolling. Accordingly, a content of Sn is 0.050% or less.
  • Sn may not necessarily be contained, and a lower limit is 0%, although the lower limit may be 0.001% as necessary.
  • the balance is Fe and impurities.
  • impurity refers to a raw material such as ore and scrap and a component contained due to various factors in production processes, and one allowed to the extent that the present invention is not adversely affected.
  • a metal microstructure in the present invention refers to a microstructure that is present at a position 1/4W or 3/4W from an end face of the steel sheet and 1/4t or 3/4t from a surface of the steel sheet.
  • a symbol "%" in the following description represents "area %”.
  • Ferrite resulting from precipitation-strengthened pro-eutectoid ferrite in which Ti carbides are finely dispersed hereafter, also referred to as "precipitation strengthened ferrite"
  • the Ti carbides being present in the ferrite due to interphase boundary precipitation or homogeneous nucleation driven by supersaturation of the Ti carbides while fine precipitates containing Ti (carbides or the like of finely precipitated Ti, and hereafter, also referred to as “fine Ti precipitates”) are subjected to ⁇ -> ⁇ transformation during cooling after rolling.
  • the precipitation strengthened ferrite is a microstructure necessary for achieving both excellent uniform elongation and strength.
  • the area fraction of the precipitation strengthened ferrite is 5 to 70%.
  • the area fraction of the precipitation strengthened ferrite is preferably 7% or more, and more preferably 10% or more.
  • the area fraction of the precipitation strengthened ferrite is preferably 65% or less, and more preferably 60% or less.
  • the precipitation strengthened ferrite of the present invention refers to ferrite in which a number density of fine Ti precipitates contained in the grains is 1.0 ⁇ 10 16 to 50.0 ⁇ 10 16 /cm 3 .
  • a number density of fine Ti precipitates contained in the grains is 1.0 ⁇ 10 16 to 50.0 ⁇ 10 16 /cm 3 .
  • the effect of precipitation strengthening cannot be obtained sufficiently.
  • the number density of fine Ti precipitates is more than 50.0 ⁇ 10 16 /cm 3 , not only the strength saturates but also the ductility decreases.
  • the condition that the area fraction of the precipitation strengthened ferrite is 5 to 70% means that the area fraction of ferrite is 5 to 70% and the number density of fine Ti precipitates contained in the ferrite grains is 1.0 ⁇ 10 16 to 50.0 ⁇ 10 16 /cm 3 .
  • an average circle-equivalent diameter of fine Ti precipitates contained in the grains of precipitation strengthened ferrite is preferably 1.00 to 3.00 nm. This is because when the average circle-equivalent diameter of fine Ti precipitates is less than 1.00 nm, the effect of precipitation strengthening is less likely to be obtained, while on the other hand, when the precipitates are coarse grained and the average circle-equivalent diameter is more than 3.00 nm, it is difficult to secure a sufficient amount of fine Ti precipitates.
  • Bainite is a microstructure necessary for balancing strength and local ductility, and has an effect of inhibiting crack propagation.
  • too much bainite leads to a decrease in ferrite, and uniform elongation is significantly degraded even though local ductility is excellent.
  • an area fraction of bainite is 30 to 95%.
  • the area fraction of bainite is preferably 80% or less, and more preferably 70% or less in the case in which uniform elongation is emphasized.
  • a high burring steel is characterized by presence of precipitation strengthened ferrite and bainite for securing workability as well as high strength such that both strength and workability are achieved.
  • presence of thermodynamically stable retained austenite, which has not been subjected to martensitic transformation, in a steel sheet indicates that the retained austenite may have high concentration of C, and hardness of martensite generated by strain induced transformation of the retained austenite during sheet metal forging may become too high, promoting void generation.
  • the amount of retained austenite is preferably as small as possible, and an area fraction of the retained austenite is 2% or less.
  • the area fraction of the retained austenite is preferably 1.5% or less, 1% or less, or 0.5% or less. It is not particularly necessary to define a lower limit, and the lower limit is 0%, which is most preferable.
  • a high burring steel is characterized by presence of precipitation strengthened ferrite and bainite for securing workability as well as high strength such that both strength and workability are achieved.
  • an area fraction of martensite which is a hard microstructure
  • a void is likely to be generated at a border between the martensite and ferrite and rupture is likely to occur as strain of a steel sheet increases by sheet metal forging.
  • an area fraction of martensite is 2% or less.
  • the area fraction of martensite is preferably 1.5% or less, 1% or less, or 0.5% or less. It is not particularly necessary to define a lower limit, and the lower limit is 0%.
  • Pearlite may serve as an initiation point of fracture during hole expansion, and thus an area fraction of pearlite is 1% or less.
  • the area fraction of pearlite is preferably 0.5% or less. It is preferable to reduce the area fraction of pearlite as much as possible and the area fraction is preferably 0%.
  • a high burring steel includes precipitation strengthened ferrite providing both excellent uniform elongation and strength and bainite providing both strength and local ductility. Excellent strength, uniform elongation, and local ductility are thereby obtained.
  • a total area fraction of precipitation strengthened ferrite and bainite is less than 95%, these properties are degraded. Accordingly, the total area fraction of precipitation strengthened ferrite and bainite is 95% or more.
  • the total area fraction is preferably 97% or more, and more preferably 98% or more.
  • an area fraction of metal microstructures is determined as follows. A sample is taken at a position 1/4W or 3/4W from an end face of the steel sheet and 1/4t or 3/4t from a surface of the steel sheet, as described above. Then, a rolling direction cross section (so-called L-direction cross section) of the sample is observed.
  • the sample is subjected to Nital etching and observed in a 300 ⁇ m ⁇ 300 ⁇ m field of view using an optical microscope after the etching. Then, a resultant microstructure photograph is subjected to image analysis to obtain an area fraction A of ferrite, an area fraction B of pearlite, and a total area fraction C of bainite, martensite, and retained austenite.
  • the portion subjected to Nital etching is subjected to Lepera etching and observed in a 300 ⁇ m ⁇ 300 ⁇ m field of view using an optical microscope. Then, a resultant microstructure photograph is subjected to image analysis to calculate a total area fraction D of retained austenite and martensite. Further, a sample subjected to facing up to a depth of 1/4 sheet thickness from a normal direction of the sheet surface is used to determine a volume ratio of the retained austenite with X-ray diffraction measurement. Since the volume ratio is substantially equal to the area fraction, the volume ratio is defined as an area fraction E of the retained austenite.
  • An area fraction of bainite is determined from a difference between the area fraction C and the area fraction D, and an area fraction of martensite is determined from a difference between the area fraction E and the area fraction D. In this way, the area fraction of each of ferrite, bainite, martensite, retained austenite, and pearlite can be determined.
  • an area fraction of precipitation strengthened ferrite can be obtained with Kernel Average Misorientation (KAM) method that comes with EBSP-OIMTM (Electron Back Scatter Diffraction Pattern-Orientation Image Microscopy).
  • KAM Kernel Average Misorientation
  • EBSP-OIMTM Electro Back Scatter Diffraction Pattern-Orientation Image Microscopy
  • KAM method averages of orientation differences among 6 adjacent pixels (first approximation) of some regular-hexagonal pixels in the measurement data, 12 pixels (second approximation) outside thereof, and 18 pixels (third approximation) further outside thereof are calculated, and each averaged value is taken as a value for the center pixel to perform calculation on each pixel.
  • a map representing an orientation variation in grains can be created by performing the calculation such that the grain boundary is not crossed. In other words, the map represents a strain distribution based on a local orientation variation in the grains.
  • an average orientation difference is calculated between adjacent pixels in the third approximation in EBSP-OIMTM, and a portion with a calculated orientation difference being 1° or less is determined to be precipitation strengthened ferrite.
  • a temperature range in which precipitation strengthened ferrite of the present invention is generated coincides with a temperature range in which interphase boundary precipitation or homogeneous nucleation of Ti carbides occurs in the ferrite by being driven by supersaturation of the Ti carbides while the Ti carbides are subjected to ⁇ -> ⁇ transformation during cooling after rolling. Since polygonal pro-eutectoid ferrite subjected to transformation at a high temperature is generated in diffusion transformation, it has a small dislocation density and less strain in grains, and thus a small inter-grain difference in crystal orientation. Accordingly, the precipitation strengthened ferrite also has a small crystal orientation difference.
  • the area fraction of precipitation strengthened ferrite was measured as follows. A sample taken in the same way as described in the microstructure observation was polished with a colloidal silica abrasive for 30 to 60 minutes, and was subjected to the EBSP measurement under the following conditions: magnification of 400x, area of 160 ⁇ m ⁇ 256 ⁇ m, and a measurement step of 0.5 ⁇ m.
  • the EBSP-OIMTM method is constituted by a device and software, with which a sample highly inclined in a scanning electron microscope (SEM) is irradiated with an electron beam, and a high sensitivity camera is used to capture a Kikuchi pattern formed by backscattering, followed by computer image processing, so that a crystal orientation at the irradiated point is measured in a short time.
  • SEM scanning electron microscope
  • the analyzable area is an area that can be observed by SEM, and the analysis can be made at a minimum resolution of 20 nm depending on the resolution of the SEM.
  • the analysis takes several hours, and a region to be analyzed is mapped to tens of thousands of points in an equidistant grid. In a polycrystalline material, a crystal orientation distribution and a size of grain in the sample can be seen.
  • a needle-shaped sample is produced from a sample to be measured by cutting and electrolytic polishing, and as necessary, with the use of focused ion beam processing along with the electrolytic polishing.
  • 3D Atom Probe Tomography accumulated data can be reconstructed to determine an actual distribution of atoms in a real space.
  • the size of the fine Ti precipitate is a circle-equivalent diameter calculated assuming that the fine Ti precipitate is spherical from the number of Ti atoms constituting the observed fine Ti precipitate and a lattice constant of the fine Ti precipitate.
  • the detection ratio ⁇ of the instrument used for measurement in the present invention was 0.35.
  • the circle-equivalent diameter is calculated as substantially 1 nm.
  • circle-equivalent diameters of 30 or more fine Ti precipitates are arbitrarily measured and an average thereof is determined.
  • the number density of fine Ti precipitates is determined with a measurement field of view being taken as a denominator and the number of the fine Ti precipitates as a numerator. In measurement of the number density, measurement was made in 5 or more fields of view, each of the field of view being 10 nm (sheet thickness direction t) ⁇ 40 nm (sheet width direction W) ⁇ 60 nm (sheet longitudinal direction L), and an average of the number densities (number of the fine Ti precipitates/cm 3 ) was determined.
  • Average circle-equivalent diameter of TiN 1.0 to 10.0 ⁇ m
  • the average circle-equivalent diameter of TiN is 10.0 ⁇ m or less.
  • the average circle-equivalent diameter of TiN is preferably 8.0 ⁇ m or less, and more preferably 5.0 ⁇ m or less.
  • the average circle-equivalent diameter of TiN is determined as follows. A sample is taken at a position 1/4W or 3/4W from an end face of the steel sheet and 1/4t or 3/4t from a surface of the steel sheet, as described above. Then, a rolling direction cross section (so-called L-direction cross section) of the sample is polished and observed with the sample unetched. Specifically, an optical microscope is used to capture a microstructure photograph at a magnification of 1000x, and the microstructure photograph is observed visually or using an image processing device and the like.
  • circle-equivalent diameters are determined for possible TiN precipitates, and those that have 1.0 ⁇ m or more of the circle-equivalent diameter are determined as TiN. Then, observation is made in 20 or more fields of view, each field of view being 60 ⁇ m(rolling direction L) ⁇ 40 ⁇ m (sheet thickness direction t), and an average of all the circle-equivalent diameters of TiN precipitates is defined as the average circle-equivalent diameter of TiN.
  • Average of minimum distances between adjacent TiN precipitates 10.0 ⁇ m or more
  • an average of distances between adjacent TiN precipitates is 10.0 ⁇ m or more.
  • the average is preferably 15.0 ⁇ m or more, and more preferably 20.0 ⁇ m or more. No upper limit is particularly defined. However, precipitation of TiN is unavoidable to some extent, and accordingly, an average of minimum distances between adjacent TiN precipitates is preferably 1000 ⁇ m or less.
  • the average of minimum distances between adjacent TiN precipitates is determined as follows. 20 TiN precipitates are arbitrarily selected, every distances between one of the TiN precipitates and another one most adjacent to it are calculated, and an average thereof is calculated. The minimum distance between TiN precipitates is measured as with the case of the average circle-equivalent diameters.
  • Standard deviation of nano hardness 1.0 GPa or less
  • voids from coalescing with each other and growing into a crack by reducing a difference in deformability between a hard microstructure and a soft microstructure to reduce voids generated at an interface between the two microstructures and to create a void spacing. Accordingly, it is possible to inhibit void generation by reducing a nano hardness difference, which corresponds to the difference in deformability between a hard microstructure and a soft microstructure.
  • a standard deviation of nano hardness in a sample cross section is employed as an indicator for a hardness difference between a soft microstructure and a hard microstructure.
  • Nano hardness can be measured with the use of, for example, TriboScope/TriboIndenter available from Hysitron.
  • the systems can arbitrarily measure nano hardness at 100 or more points at a load of 1 mN, and calculate a standard deviation of the nano hardness from the results.
  • a smaller standard deviation of nano hardness is preferable, and accordingly, it is 1.0 GPa or less.
  • the standard deviation of nano hardness is preferably 0.8 GPa or less.
  • the steel sheet according to the present invention preferably has a tensile strength of 780 MPa or more, which is a similar level to a conventional high burring steel. It is not particularly necessary to define an upper limit to the tensile strength. However, it may be 1200 MPa, 1150 MPa, or 1000 MPa.
  • the tensile strength refers to a tensile strength in JIS Z 2241 (2011).
  • a small uniform elongation is likely to be a cause of sheet thickness reduction due to necking during press forming, and then a cause of press cracking.
  • uniform elongation u ⁇ EL ln ⁇ n 0 + 1 where in a relation between a nominal stress ⁇ n and a nominal strain ⁇ n, ⁇ n0 is a nominal strain at a point where a value obtained by differentiating the nominal stress ⁇ n with the nominal strain ⁇ n is zero.
  • a poor hole expandability leads to a poor material flowability during stretch flanging, and may possibly cause a cracking.
  • the hole expansion ratio ( ⁇ ) represents a hole expansion ratio ( ⁇ ) according to a test method in conformity with JIS Z 2256 (2010).
  • the equivalent plastic strain is converted using a relation between a shearing stress ⁇ s and a shear plastic strain ⁇ sp in a simple shear test into a relation between a tensile stress ⁇ and a tensile strain ⁇ in a uniaxial tensile test, which is different in deformation mode, and a constant, conversion factor ( ⁇ ) is used to make a conversion, assuming an isotropic hardening rule and a plastic work conjugate relationship.
  • the isotropic hardening rule is a work hardening rule in which it is assumed that the shape of yield curve does not change even when a strain develops (that is, it expands in a similar shape).
  • the plastic work conjugate relationship is a relationship in which work hardening is described only as a function of a plastic work, and exhibits the same amount of work hardening given the same plastic work ( ⁇ ⁇ ⁇ ) regardless of the deformation mode.
  • a shearing stress and a shear plastic strain in a simple shear test can thereby converted into a tensile stress and a tensile strain in a uniaxial tensile test.
  • the relation is shown below.
  • conversion factor ⁇ is determined such that a relation between a shearing stress and a shear plastic strain is similar to a relation between a tensile stress and a tensile strain.
  • the conversion factor ⁇ can be determined in the following procedure. First, a relation between a tensile strain ⁇ (actual value) and a tensile stress ⁇ (actual value) in a uniaxial tensile test is determined. Then, a relation between a shearing strain ⁇ s (actual value) and a shearing stress ⁇ s (actual value) in a uniaxial shear test.
  • ⁇ is changed to determine a tensile strain ⁇ (converted) determined from the shearing strain ⁇ s (actual value) and a tensile stress ⁇ (converted) determined from the shearing stress ⁇ s (actual value). Then, the tensile stress ⁇ (converted) when the tensile strain ⁇ (converted) is from 0.2% to uniform elongation (u-EL) is determined. At this time, an error between the tensile stress ⁇ (converted) and the tensile stress ⁇ (actual value) is determined, and " ⁇ " that minimizes the error is determined with the method of least squares.
  • An equivalent plastic strain ⁇ eq is defined as a shear plastic strain ⁇ sp (rupture) at the time of rupture in a simple shear test converted, with the use of the determined ⁇ , into a tensile strain ⁇ in a simple tensile test.
  • the steel sheet according to the present invention is characterized by good workability in a high strain domain typified by sheet metal forging, and its equivalent plastic strain ⁇ eq satisfies 0.50 or more. Since the equivalent plastic strain of a conventional TRIP steel at best on the order of 0.30, it has been confirmed that the steel sheet according to the present invention has a good sheet forgeability.
  • the steel sheet according to the present invention finds application primarily in automobiles and the like and the sheet thickness is ranging primarily from 1.0 to 4.0 mm. Accordingly, the range of sheet thickness may be from 1.0 to 4.0 mm, and, as necessary, a lower limit may be 1.2 mm, 1.4 mm, or 1.6 mm, and an upper limit may be 3.6 mm, 3.2 mm, or 2.8 mm.
  • the present inventors confirmed that the hot rolled steel sheet of the present invention can be produced by conducting the following production processes (a) to (f) in this order. Description will now be made as to each of the production processes in detail.
  • Production methods prior to hot rolling are not particularly limited except a deoxidation method for molten steel.
  • sequential deoxidation described below can be executed for controlling a composition and size of TiN.
  • a variety of second smelting including the sequential deoxidation is executed to make an adjustment for a component composition described above.
  • methods such as general continuous casting and thin slab casting may be used to produce a slab.
  • scrap or the like may be used as raw materials provided that the material can be controlled into the component range of the present invention.
  • a dissolved oxygen concentration in the molten steel is first adjusted with the amount of Si, which is a weaker deoxidizing element than Ti, to bring the dissolved oxygen concentration in equilibrium with the amount of Si into 0.0020 to 0.0080%.
  • the dissolved oxygen concentration is less than 0.0020%, the amount of complex oxides enough to eventually make the size of nitrides including Ti smaller cannot be obtained.
  • the dissolved oxygen concentration is more than 0.0080%, generated complex oxides are coarsened, so that the effect of making the size of nitrides including Ti smaller is impaired.
  • the content of Si is less than 0.05%, the dissolved oxygen concentration in equilibrium with Si reaches more than 0.0080%.
  • the content of Si is more than 0.2%, the dissolved oxygen concentration in equilibrium with Si falls less than 0.0020%. Accordingly, the content of Si at a stage preceding to deoxidation is 0.05 to 0.20%, and the dissolved oxygen concentration is 0.0020 to 0.0080%.
  • Ti is added such that a final content of Ti is 0.015 to 0.170%
  • Al is added such that a final content of A1 is 0.005 to 0.020% to execute the deoxidation.
  • Ti oxides generated when Ti is added grow and are coagulated and coarsened, resulting in flotation. Accordingly, addition of A1 is desirably done immediately after Ti is added.
  • deoxidation is executed by adding Ca, which is an even stronger deoxidizing element than Ti and Al.
  • Ca may be added such that a final content of Ca is 0.0005 to 0.0100%.
  • addition of Ca is in such an amount that the final content of Ca falls less than 0.0005%, the amounts of Ti, Al, Ca, or oxides thereof obtained in the sequential deoxidation are reduced, and thus the amount of TiN precipitated with those elements being as cores is reduced. Accordingly, a high-strength hot-rolled steel sheet excellent for sheet metal forging cannot be obtained.
  • Mg or REM which are stronger deoxidizing elements than Ti and Al, may be used for deoxidation instead of Ca. Specifically, within 5 minutes after Al is added, deoxidation is executed by adding Mg, which is an even stronger deoxidizing element than Ti and Al, to the extent that a final content of Mg is 0.0005 to 0.0100%. Similarly, within 5 minutes after Al is added, deoxidation may be executed by adding REM, which is an even stronger deoxidizing element than Ti and Al, to the extent that a final content of REM is 0.0005 to 0.1000%.
  • two or more selected from Ca, Mg, and REM may be added in combination.
  • one or more selected from Ca, Mg, and REM may be added such that a final total content of Ca, Mg, and REM is 0.0005 or more, a final content of Ca and Mg is 0.0100% or less, and a final content of REM is 0.1000% or less.
  • a produced slab is subjected to hot rolling into a hot rolled steel sheet.
  • the slab is heated first.
  • the slab is heated to a temperature of SRTmin°C or more represented by the following formula (i), or to 1260°C or less.
  • SRTmin refers to a solution treatment temperature of TiC.
  • SRTmin 7000 / 2.75 ⁇ log Ti ⁇ C ⁇ 273 where a symbol of an element in the above formula represents a content (in mass %) of the element in the hot rolled steel sheet and is substituted by zero when the element is not contained.
  • the slab extracted from a heating furnace is subjected to rough rolling and subsequent multi-stand finish rolling.
  • an end temperature of the rough rolling is 1100°C or more.
  • the multi-stand finish rolling is conducted by continuous rolling at multiple, three stands or more (for example, 6 or 7 stands). The multi-stand finish rolling is executed such that a cumulative strain (effective cumulative strain) of rolling at final three stands is 0.01 to 0.10.
  • the effective cumulative strain is an indicator that takes into consideration a grain size variation according to temperature during rolling and rolling reduction of a steel sheet by rolling and a grain size variation when grains statically recover in a time lapse after rolling.
  • An end temperature of the multi-stand finish rolling may be satisfactory if it is Ar 3 (°C) + 30°C or more where Ar 3 is determined in the following formula (ii). This is because precipitation strengthened ferrite and bainite intended in the present invention can thus be obtained.
  • Ar 3 970 ⁇ 325 ⁇ C + 33 ⁇ Si + 287 ⁇ P + 40 ⁇ Al ⁇ 92 ⁇ Mn + Mo + Cu ⁇ 46 ⁇ Cr + Ni where a symbol of an element in the above formula represents a content (in mass %) of the element in the hot rolled steel sheet and is substituted by zero when the element is not contained.
  • cooling of the resultant hot rolled steel sheet is started after 1.00 to 5.00 seconds. Then, the sheet is cooled at an average cooling rate of 10°C/sec or more from the rolling end temperature down to a temperature of 650 to 750°C, and thereafter the sheet is held in air for 1 to 10 seconds.
  • the average cooling rate of the first cooling process is less than 10°C/sec, pearlite is likely to be generated.
  • an upper limit of the cooling rate is not particularly limited. However, it may be 300°C/sec or less in order to avoid supercooling.
  • a holding temperature in air is less than 650°C, bainite is likely to be generated and the bainite area fraction increases.
  • a holding temperature in air is more than 750°C, pearlite is likely to be generated.
  • holding in air includes a hot-rolled steel sheet being subjected to air cooling or minimally limited cooling in a cooling facility, and a lower limit at this time is ideally 0°C/sec and an upper limit is 8°C/sec.
  • cooling is conducted at an average cooling rate of 10°C/sec or more from a temperature range of 600 to 740°C.
  • a cooling starting temperature less than 600°C prevents ferritic transformation from sufficiently progressing and causes insufficient precipitation of fine Ti precipitates.
  • ferritic transformation may excessively progress and pearlite may be generated, leading to degradation of hole expandability.
  • fine Ti precipitates may be coarsened to reduce strength.
  • the average cooling rate is less than 10°C/sec, pearlite may also be generated, leading to degradation of hole expandability.
  • An upper limit of the average cooling rate is not particularly limited. However, it may be 1000°C/sec or less because there is a concern that the steel sheet warps due to thermal strain caused by thermal deviation.
  • the cooled hot rolled steel sheet is coiled at a coiling temperature of 450 to 650°C.
  • Conditions after coiling process are not particularly limited.
  • the hot-rolled steel sheet thus obtained has excellent sheet forgeability. Accordingly, the hot-rolled steel sheet can be forged in sheet metal forging and the like to obtain forged parts in complex shapes, which are required to have high strength that could not have been achieved in the past.
  • the present inventors observed metal microstructures of the resultant hot rolled steel sheet and measured the area fraction of each of the microstructures. Specifically, when a width and a thickness of the steel sheet in a cross section perpendicular to a rolling direction of the steel sheet are defined as W and t, respectively, a specimen for metal microstructure observation was cut out at a position 1/4W from an end face of the steel sheet and 1/4t from a surface of the steel sheet.
  • a rolling direction cross section (so-called L-direction cross section) of the specimen was subjected to Nital etching, and observed in a 300 ⁇ m ⁇ 300 ⁇ m field of view using an optical microscope after the etching. Then, a resultant microstructure photograph was subjected to image analysis to determine an area fraction A of ferrite, an area fraction B of pearlite, and a total area fraction C of bainite, martensite, and retained austenite.
  • the portion subjected to Nital etching was subjected to Lepera etching and observed in a 300 ⁇ m ⁇ 300 ⁇ m field of view using an optical microscope. Then, a resultant microstructure photograph was subjected to image analysis to calculate a total area fraction D of retained austenite and martensite. Further, a sample subjected to facing up to a depth of 1/4 sheet thickness from a normal direction of the sheet surface was used to determine a volume ratio of the retained austenite with X-ray diffraction measurement. Since the volume ratio is substantially equal to the area fraction, the volume ratio was defined as an area fraction E of the retained austenite.
  • An area fraction of bainite was determined from a difference between the area fraction C and the area fraction D, and an area fraction of martensite was determined from a difference between the area fraction E and the area fraction D. In this way, the area fraction of each of ferrite, bainite, martensite, retained austenite, and pearlite was determined.
  • the area fraction of precipitation strengthened ferrite was determined in KAM method after the specimen was polished with a colloidal silica abrasive and subjected to the EBSP measurement under the following conditions: magnification of 400x, 160 ⁇ 256 ⁇ m field of view, and a measurement step of 0.5 ⁇ m, as described above.
  • the specimen was subjected to electrolytic polishing and measured in the 3D Atom Probe Tomography to determine the circle-equivalent diameter and the number density, as described above.
  • the specimen was observed at a magnification of 1000x in 20 fields of view, each field of view being 60 ⁇ 40 ⁇ m, and the average circle-equivalent diameter of TiN was determined by image processing, as described above. Further, the minimum distance between TiN precipitates was determined by observing the same portions as in the microstructure studies using a metallurgical microscope at a magnification of 500x.
  • tensile strength properties tensile strength (TS), uniform elongation (u-EL), and hole expansion ratio ( ⁇ )
  • TS tensile strength
  • u-EL uniform elongation
  • ⁇ hole expansion ratio
  • JIS Z 2241 (2011) No. 5 specimen, which was taken at a position 1/4W or 3/4W from one end of the sheet in a sheet width direction when a sheet width is defined as W with a direction (width direction) perpendicular to a rolling direction being a longitudinal direction.
  • the hole expansion ratio was evaluated in conformity with a test method described in JIS Z 2256 2010 using a specimen taken at a position similar to the position where the tensile test specimen was taken.
  • the present inventors conducted a simple shear test in a procedure described below, and determined the equivalent plastic strain based on the results.
  • a specimen for the simple shear test is taken at a position 1/4W or 3/4W from one end of the sheet in a sheet width direction when a sheet width is defined as W with a direction (width direction) perpendicular to a rolling direction being a longitudinal direction.
  • Figure 1(a) illustrates an example of the specimen.
  • the specimen for the simple shear test illustrated in Figure 1(a) was processed into a rectangular specimen of 23 mm in the width direction of the steel sheet and 38 mm in the rolling direction of the steel sheet in such a way that both sides were uniformly polished to a sheet thickness of 2.0 mm for uniform sheet thickness.
  • Chucks were applied to opposite chucking portions 2 on long sides (rolling direction) of the specimen, each chucking portion having 10 mm along a short side direction (width direction), so that a shear width (shear deformation generation portion 1) of 3 mm is provided in the center of the specimen.
  • the sheet thickness is less than 2.0 mm
  • the test was conducted with the sheet thickness being left intact without polishing. Further, the center of the specimen was marked with a straight line in the short side direction (width direction) with a pen or the like.
  • FIG. 1(b) illustrates an example of the specimen subjected to shear deformation.
  • shear nominal stress is nearly equal to the shear true stress.
  • the specimen for the metal microstructure observation was polished again.
  • the specimen was measured in measurement areas of 25 ⁇ m ⁇ 25 ⁇ m each at an interval of 5 ⁇ m at a 1/4 depth position (1/4t portion) of sheet thickness t from a steel sheet surface in a cross section in parallel to the rolling direction under a load of 1 mN (loading 10s and unloading 10s). From the results, an average nano hardness value and a standard deviation of nano hardness were calculated.
  • the nano hardness was measured with the use of TriboScope/TriboIndenter available from Hysitron.
  • a hot-rolled steel sheet having balanced properties which has a tensile strength (TS) of 780 MPa or more, a product (TS ⁇ u-EL) of a uniform elongation u-EL and the tensile strength TS being equal to 7000 MPa ⁇ % or more, and a product (TS ⁇ ⁇ ) of a hole expansion ratio ⁇ and the tensile strength TS being equal to 50000 MPa ⁇ % or more, can be obtained.
  • the hot rolled steel sheet according to the present invention has an equivalent plastic strain of more than 0.90 (90%), and it has been confirmed that the steel sheet can endure in high strain range working such as sheet metal forging.
  • a hot rolled steel sheet with excellent sheet forgeability maintaining a good hole expandability which is a basic feature as a high burring steel, can be obtained. Accordingly, the hot rolled steel sheet according to the present invention can find broad application in machine parts and the like. In particular, when it is applied to working on steel sheets including working in a high strain range such as sheet metal forging, remarkable effects thereof can be achieved.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Mechanical Engineering (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Heat Treatment Of Sheet Steel (AREA)
  • Heat Treatment Of Steel (AREA)
  • Treatment Of Steel In Its Molten State (AREA)
EP17904326.0A 2017-03-31 2017-03-31 Warmgewalztes stahlblech, stahlblechteil und herstellungsverfahren dafür Withdrawn EP3604587A4 (de)

Applications Claiming Priority (1)

Application Number Priority Date Filing Date Title
PCT/JP2017/013749 WO2018179391A1 (ja) 2017-03-31 2017-03-31 熱間圧延鋼板および鋼製鍛造部品ならびにそれらの製造方法

Publications (2)

Publication Number Publication Date
EP3604587A1 true EP3604587A1 (de) 2020-02-05
EP3604587A4 EP3604587A4 (de) 2020-09-09

Family

ID=62236389

Family Applications (1)

Application Number Title Priority Date Filing Date
EP17904326.0A Withdrawn EP3604587A4 (de) 2017-03-31 2017-03-31 Warmgewalztes stahlblech, stahlblechteil und herstellungsverfahren dafür

Country Status (8)

Country Link
US (1) US20200024679A1 (de)
EP (1) EP3604587A4 (de)
JP (1) JP6332571B1 (de)
KR (1) KR20190126100A (de)
CN (1) CN110475890A (de)
BR (1) BR112019017622A2 (de)
MX (1) MX2019011711A (de)
WO (1) WO2018179391A1 (de)

Families Citing this family (14)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN111801436B (zh) * 2019-02-05 2021-10-29 日本制铁株式会社 钢构件、钢板及它们的制造方法
JP7417171B2 (ja) * 2020-09-30 2024-01-18 日本製鉄株式会社 鋼材
EP4234742A4 (de) * 2021-01-12 2025-04-30 Nippon Steel Corporation Warmgewalztes stahlblech
CN112981237B (zh) * 2021-01-28 2022-10-11 江阴兴澄特种钢铁有限公司 一种球笼式万向节保持架用钢及其生产方法
DE102021104584A1 (de) * 2021-02-25 2022-08-25 Salzgitter Flachstahl Gmbh Hochfestes, warmgewalztes Stahlflachprodukt mit hoher lokaler Kaltumformbarkeit sowie ein Verfahren zur Herstellung eines solchen Stahlflachprodukts
WO2023008163A1 (ja) * 2021-07-27 2023-02-02 日本製鉄株式会社 鋼板およびその製造方法
EP4400613A4 (de) * 2021-09-06 2024-11-06 Nippon Steel Corporation Warmgewalztes stahlblech
KR20250057878A (ko) * 2022-10-26 2025-04-29 제이에프이 스틸 가부시키가이샤 고강도 강판, 부재 및 그들의 제조 방법
JP7367893B1 (ja) * 2022-12-08 2023-10-24 Ｊｆｅスチール株式会社 高強度鋼板、高強度鋼板を用いてなる部材、部材からなる自動車の骨格構造部品用又は自動車の補強部品、ならびに高強度鋼板及び部材の製造方法
CN116230143B (zh) * 2023-04-27 2023-07-11 燕山大学 一种提升变厚度金属板带材延伸率的设计方法
CN117165848A (zh) * 2023-08-15 2023-12-05 中国科学院金属研究所 一种耐海洋大气腐蚀及缝隙腐蚀的螺栓合金及其制备方法
CN118166288A (zh) * 2024-01-10 2024-06-11 湖南华菱涟源钢铁有限公司 高强耐候钢及其生产方法
CN117947354A (zh) * 2024-01-31 2024-04-30 湖南华菱涟源钢铁有限公司 耐候热轧钢带及其生产方法、耐候钢带
CN119464656B (zh) * 2024-11-25 2025-07-11 中北大学 一种改善特超级铁素体不锈钢热轧边裂的锻轧方法

Family Cites Families (9)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP3790135B2 (ja) 2000-07-24 2006-06-28 株式会社神戸製鋼所 伸びフランジ性に優れた高強度熱延鋼板およびその製造方法
JP3888128B2 (ja) 2000-10-31 2007-02-28 Ｊｆｅスチール株式会社 材質均一性に優れた高成形性高張力熱延鋼板ならびにその製造方法および加工方法
JP3882577B2 (ja) 2000-10-31 2007-02-21 Ｊｆｅスチール株式会社 伸びおよび伸びフランジ性に優れた高張力熱延鋼板ならびにその製造方法および加工方法
JP4516924B2 (ja) * 2006-03-23 2010-08-04 新日本製鐵株式会社 熱間圧延時の耐表面割れ性に優れた薄鋼板及びその製造方法
JP5194857B2 (ja) * 2008-02-08 2013-05-08 Ｊｆｅスチール株式会社 高強度熱延鋼板およびその製造方法
WO2012127125A1 (fr) * 2011-03-24 2012-09-27 Arcelormittal Investigatión Y Desarrollo Sl Tôle d'acier laminée à chaud et procédé de fabrication associé
JP6194951B2 (ja) * 2013-04-15 2017-09-13 新日鐵住金株式会社 熱延鋼板
JP6036617B2 (ja) * 2013-09-10 2016-11-30 Ｊｆｅスチール株式会社 靭性に優れた高強度熱延鋼板およびその製造方法
KR101536471B1 (ko) * 2013-12-24 2015-07-13 주식회사 포스코 용접열영향부 인성이 우수한 초고강도 용접구조용 강재 및 이의 제조방법

Also Published As

Publication number Publication date
WO2018179391A1 (ja) 2018-10-04
JP6332571B1 (ja) 2018-05-30
BR112019017622A2 (pt) 2020-03-24
KR20190126100A (ko) 2019-11-08
CN110475890A (zh) 2019-11-19
EP3604587A4 (de) 2020-09-09
JPWO2018179391A1 (ja) 2019-04-11
MX2019011711A (es) 2019-11-21
US20200024679A1 (en) 2020-01-23

Similar Documents

Publication Publication Date Title
EP3604583A1 (de) Warmgewalztes stahlblech, stahlblechteil und herstellungsverfahren dafür
EP3604587A1 (de) Warmgewalztes stahlblech, stahlblechteil und herstellungsverfahren dafür
EP3282031B1 (de) Wärmebehandeltes stahlblechelement und herstellungsverfahren dafür
US8876986B2 (en) Cold-rolled steel sheet
EP3604586A1 (de) Warmgewalztes stahlblech
EP2251448B1 (de) Kaltgewalzte stahlbleche
EP2823905B1 (de) Heisspressformverfahren und karosseriekomponenten eines kraftfahrzeugs
KR101617115B1 (ko) 열연 강판 및 그 제조 방법
EP3282030B1 (de) Wärmebehandeltes stahlblechelement und herstellungsverfahren dafür
EP3604585A1 (de) Warmgewalztes stahlblech
CN101861406B (zh) 高强度冷轧钢板
EP3495529B1 (de) Stahlblech und plattiertes stahlblech
EP3495527A1 (de) Stahlblech und plattiertes stahlblech
EP2039791A1 (de) Hochfestes stahlblech und herstellungsverfahren dafür
EP3868909A1 (de) Dünnes stahlblech und verfahren zur herstellung davon
EP4095272A1 (de) Stahlblech und verfahren zur herstellung davon
KR20170107057A (ko) 고강도 냉연 강판 및 그의 제조 방법
EP3460088B1 (de) Stahlblech
TWI635187B (zh) Hot rolled steel sheet and steel forged parts and manufacturing method thereof
JP7836012B2 (ja) 鋼板及びその製造方法
TWI639713B (zh) Hot rolled steel sheet and steel forged parts and manufacturing method thereof
EP4682282A1 (de) Stahlplatte

Legal Events

Date Code Title Description
STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: THE INTERNATIONAL PUBLICATION HAS BEEN MADE

PUAI Public reference made under article 153(3) epc to a published international application that has entered the european phase

Free format text: ORIGINAL CODE: 0009012

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: REQUEST FOR EXAMINATION WAS MADE

17P Request for examination filed

Effective date: 20190828

AK Designated contracting states

Kind code of ref document: A1

Designated state(s): AL AT BE BG CH CY CZ DE DK EE ES FI FR GB GR HR HU IE IS IT LI LT LU LV MC MK MT NL NO PL PT RO RS SE SI SK SM TR

AX Request for extension of the european patent

Extension state: BA ME

DAV Request for validation of the european patent (deleted)
DAX Request for extension of the european patent (deleted)
A4 Supplementary search report drawn up and despatched

Effective date: 20200810

RIC1 Information provided on ipc code assigned before grant

Ipc: C21D 9/48 20060101ALI20200804BHEP

Ipc: C21D 9/46 20060101ALI20200804BHEP

Ipc: C21D 8/04 20060101ALI20200804BHEP

Ipc: C22C 38/14 20060101ALI20200804BHEP

Ipc: C21D 1/02 20060101ALI20200804BHEP

Ipc: C21C 7/06 20060101ALI20200804BHEP

Ipc: C21D 8/02 20060101ALI20200804BHEP

Ipc: C22C 38/04 20060101ALI20200804BHEP

Ipc: C21D 6/00 20060101ALI20200804BHEP

Ipc: C22C 38/58 20060101ALI20200804BHEP

Ipc: C21D 1/19 20060101ALI20200804BHEP

Ipc: C21C 7/04 20060101ALI20200804BHEP

Ipc: C22C 38/02 20060101ALI20200804BHEP

Ipc: C22C 38/00 20060101AFI20200804BHEP

Ipc: C22C 38/06 20060101ALI20200804BHEP

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: THE APPLICATION IS DEEMED TO BE WITHDRAWN

18D Application deemed to be withdrawn

Effective date: 20210310