[go: up one dir, main page]
More Web Proxy on the site http://driver.im/

EP3221483B1 - Acier multiphases autotrempant en profondeur à haute résistance mécanique avec excellent aptitude au faconnage et procédé de fabrication d'une bande de cet acier - Google Patents

Acier multiphases autotrempant en profondeur à haute résistance mécanique avec excellent aptitude au faconnage et procédé de fabrication d'une bande de cet acier Download PDF

Info

Publication number
EP3221483B1
EP3221483B1 EP15821018.7A EP15821018A EP3221483B1 EP 3221483 B1 EP3221483 B1 EP 3221483B1 EP 15821018 A EP15821018 A EP 15821018A EP 3221483 B1 EP3221483 B1 EP 3221483B1
Authority
EP
European Patent Office
Prior art keywords
steel strip
steel
hot
content
strip
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.)
Active
Application number
EP15821018.7A
Other languages
German (de)
English (en)
Other versions
EP3221483A1 (fr
Inventor
Thomas Schulz
Joachim SCHÖTTLER
Sascha KLUGE
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.)
Salzgitter Flachstahl GmbH
Original Assignee
Salzgitter Flachstahl GmbH
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 Salzgitter Flachstahl GmbH filed Critical Salzgitter Flachstahl GmbH
Publication of EP3221483A1 publication Critical patent/EP3221483A1/fr
Application granted granted Critical
Publication of EP3221483B1 publication Critical patent/EP3221483B1/fr
Active legal-status Critical Current
Anticipated expiration legal-status Critical

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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/52—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length
    • 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
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • 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
    • C21D1/84—Controlled slow cooling
    • 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
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • 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
    • C21D6/007—Heat treatment of ferrous alloys containing Co
    • 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties 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 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 by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties 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
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0278—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving a particular surface treatment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • 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
    • 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
    • C21D9/52—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length
    • C21D9/54—Furnaces for treating strips or wire
    • C21D9/56—Continuous furnaces for strip or wire
    • C21D9/561—Continuous furnaces for strip or wire with a controlled atmosphere or vacuum
    • 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
    • C21D9/52—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length
    • C21D9/54—Furnaces for treating strips or wire
    • C21D9/56—Continuous furnaces for strip or wire
    • C21D9/562—Details
    • 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
    • C21D9/52—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length
    • C21D9/54—Furnaces for treating strips or wire
    • C21D9/56—Continuous furnaces for strip or wire
    • C21D9/58—Continuous furnaces for strip or wire with heating by baths
    • 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
    • C22C38/02—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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • 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
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • 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
    • 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
    • C23C2/003—Apparatus
    • C23C2/0038—Apparatus characterised by the pre-treatment chambers located immediately upstream of the bath or occurring locally before the dipping process
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
    • C23C2/0224—Two or more thermal pretreatments
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C23C2/024—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by cleaning or etching
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/26—After-treatment
    • C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
    • C23C2/29—Cooling or quenching
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/34—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the shape of the material to be treated
    • C23C2/36—Elongated material
    • C23C2/40—Plates; Strips
    • 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
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of annealing
    • C21D1/28—Normalising
    • 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
    • C21D1/74—Methods of treatment in inert gas, controlled atmosphere, vacuum or pulverulent material
    • 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
    • C21D1/74—Methods of treatment in inert gas, controlled atmosphere, vacuum or pulverulent material
    • C21D1/76—Adjusting the composition of the atmosphere
    • 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
    • C21D2241/00—Treatments in a special environment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0447—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0447—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the heat treatment
    • C21D8/0473—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

Definitions

  • the invention relates to a method for producing a cold-rolled or hot-rolled steel strip from an air-hardenable multiphase steel according to claim 1, and advantageous developments according to claims 2 to 20.
  • the invention relates to steels with a tensile strength in the range of at least 950 MPa in the non-tempered state for the production of components which have improved formability (such as increased hole expansion and increased bending angle) and improved welding properties.
  • a tempering treatment of these steels according to the invention can increase the yield strength and tensile strength, for example by air hardening with optional subsequent tempering.
  • the weight of the vehicles can be reduced with improved forming and component behavior during production and operation.
  • High-strength to ultra-high-strength steels must therefore have comparatively high requirements with regard to their strength and ductility, energy absorption and processing, such as punching, hot and cold forming, thermal hardening (e.g. air hardening, press hardening), welding and / or surface treatment, e.g. metallic finishing, organic coating or painting are sufficient.
  • energy absorption and processing such as punching, hot and cold forming, thermal hardening (e.g. air hardening, press hardening), welding and / or surface treatment, e.g. metallic finishing, organic coating or painting are sufficient.
  • Newly developed steels must therefore face the increasing weight requirements due to reduced sheet thickness, the increasing material requirements for yield strength, tensile strength, strengthening behavior and elongation at break with good processing properties such as formability and weldability.
  • high-strength, high-strength steel with a single-phase or multi-phase structure must be used to ensure sufficient strength of the motor vehicle components and to meet the high component requirements with regard to toughness, edge crack resistance, improved bending angle and bending radius, energy absorption as well as strengthening and bake hardening Effect to suffice.
  • the hole expansion capacity is a material property that describes the resistance of the material to crack initiation and crack propagation during forming operations in areas close to edges, such as when pulling a collar.
  • the hole expansion test is regulated, for example, in ISO 16630. Then prefabricated holes punched into a sheet, for example, are expanded using a mandrel. The measured variable is the change in the hole diameter in relation to the initial diameter at which the first crack through the sheet occurs at the edge of the hole.
  • Improved edge crack resistance means an increased formability of the sheet edges and can be described by an increased hole expansion capacity. This situation is known under the synonyms “L ow E dge C rack” (LEC) or under “H igh E H ole xpansion” (HHE) and xpand®.
  • the bending angle describes a material property that gives conclusions about the material behavior during forming operations with dominant bending components (e.g. when folding) or also in the event of crash loads. Increased bending angles thus increase passenger compartment safety.
  • the determination of the bending angle ( ⁇ ) is e.g. normatively regulated via the plate bending test in VDA 238-100.
  • the above-mentioned properties are important for components that e.g. can be formed into very complex components by air hardening with optional tempering.
  • High-strength components must be sufficiently resistant to embrittlement of hydrogen.
  • Test of resistance of A dvanced H igh S trength S Teels (AHSS) for the automotive industry with respect to hydrogen-induced production-related brittle fractures is regulated in the SEP1970 and tested on the sample and the bracket Lochzugprobe.
  • Dual-phase steels are increasingly being used in vehicle construction, which consist of a ferritic basic structure in which a martensitic second phase is embedded. It has been found that, in the case of low-carbon, micro-alloyed steels, portions of further phases such as bainite and residual austenite are advantageous, for example on the hole expansion behavior, the bending behavior and the affect hydrogen-induced brittle fracture behavior.
  • the bainite can be present in different forms, such as upper and lower bainite.
  • Multi-phase steels include e.g. Complex phase steels, ferritic-bainitic steels, TRIP steels, as well as the previously described dual phase steels, which are characterized by different structural compositions.
  • these complex phase steels Compared to dual-phase steels, these complex phase steels have higher yield strengths, a greater yield strength ratio, less strain hardening and a higher hole expansion capacity.
  • ferritic-bainitic steels are steels which contain bainite or solidified bainite in a matrix of ferrite and / or solidified ferrite.
  • the strength of the matrix is brought about by a high dislocation density, grain refinement and the excretion of microalloying elements.
  • dual-phase steels are steels with a ferritic basic structure in which a martensitic second phase is embedded in the form of an island, sometimes also with parts of bainite as the second phase. With high tensile strength, dual-phase steels show a low yield ratio and strong strain hardening.
  • TRIP steels are, according to EN 10346, steels with a predominantly ferritic structure, in which bainite and residual austenite is embedded, which can convert to martensite during the forming (TRIP effect). Due to its strong strain hardening, the steel achieves high values of uniform elongation and tensile strength. In connection with the bake hardening effect high component strengths can be achieved. These steels are suitable for both stretch drawing and deep drawing. However, higher sheet metal holder forces and press forces are required for material forming. A comparatively strong springback must be taken into account.
  • the high-strength steels with a single-phase structure include e.g. bainitic and martensitic steels.
  • Bainitic steels are, according to EN 10346, steels which are characterized by a very high yield strength and tensile strength with sufficient elongation for cold forming processes. Due to the chemical composition, it is easy to weld.
  • the structure typically consists of bainite.
  • the structure may occasionally contain small amounts of other phases, such as martensite and ferrite.
  • martensitic steels are steels that contain small amounts of ferrite and / or bainite in a basic structure of martensite through thermomechanical rolling. This steel grade is characterized by a very high yield strength and tensile strength with sufficient elongation for cold forming processes. Within the group of multi-phase steels, the martensitic steels have the highest tensile strength values. The suitability for deep drawing is limited. The martensitic steels are primarily suitable for bending forming processes, such as roll forming.
  • High and high-strength multi-phase steels are used, among others. in structural, chassis and crash-relevant components, as sheet metal blanks, tailored blanks (welded blanks) and as flexible cold-rolled strips, so-called TRB®s or tailored strips.
  • T ailor R olled B lank lightweight technology enables a significant weight reduction through a load-adjusted sheet thickness over the component length and / or steel grade.
  • a special heat treatment takes place for the defined structure adjustment, where e.g. due to comparatively soft components such as ferrite or bainitic ferrite, the steel has its low yield strength and due to its hard components such as martensite or carbon-rich bainite, its strength.
  • cold-rolled high-strength steel strips are usually recrystallized in a continuous annealing process to form sheet metal that is easy to form.
  • the process parameters such as throughput speed, annealing temperatures and cooling speed (cooling gradients), are set according to the required mechanical-technological properties with the necessary structure.
  • the pickled hot strip in typical thicknesses between 1.50 to 4.00 mm or cold strip in typical thicknesses from 0.50 to 3.00 mm is heated to a temperature in the continuous annealing furnace that during recrystallization and cooling sets the required structure formation.
  • a constant temperature is difficult to achieve, especially with different thicknesses in the transition area from one belt to another belt.
  • this can lead to e.g. the thinner strip is either passed through the furnace too slowly, which reduces productivity, or the thicker strip is passed through the furnace too quickly and the necessary annealing temperatures and cooling gradients are not achieved to achieve the desired structure.
  • the consequences are increased waste and high costs of errors.
  • TRB®s with a multi-phase structure is not without additional effort with today's known alloys and available continuous annealing systems for widely varying strip thicknesses, e.g. an additional heat treatment before cold rolling (hot strip soft annealing).
  • strip thicknesses e.g. an additional heat treatment before cold rolling (hot strip soft annealing).
  • hot strip soft annealing e.g. an additional heat treatment before cold rolling
  • a homogeneous multi-phase structure cannot be set in cold as well as hot-rolled steel strips due to a temperature gradient occurring in the usual alloy-specific narrow process windows.
  • a method for producing a steel strip with different thickness over the length of the strip is described, for example, in DE 100 37 867 A1 described.
  • the annealing treatment is usually carried out in a continuous annealing furnace upstream of the hot-dip bath.
  • the required structure is only set during the annealing treatment in the continuous annealing furnace in order to achieve the required mechanical properties.
  • Decisive process parameters are therefore the setting of the annealing temperature and the speed, as well as the cooling rate (cooling gradient) in continuous annealing, since the phase change takes place depending on the temperature and time.
  • the areas with a smaller strip thickness due to the conversion processes during cooling either have too high strengths due to excessively high martensite contents, or the areas with greater strip thickness achieve insufficient strengths due to insufficiently low martensite contents due to the process window being too small. Homogeneous mechanical-technological properties across the strip length or width can practically not be achieved with the known alloy concepts for continuous annealing.
  • the goal of achieving the resulting mechanical-technological properties in a narrow range across the bandwidth and strip length through the controlled adjustment of the volume fractions of the structural components has top priority and is only possible through an enlarged process window.
  • the known alloy concepts are characterized by an excessively narrow process window and are therefore unsuitable for solving the present problem, particularly in the case of flexibly rolled strips. With the known alloy concepts, only steels of a strength class with defined cross-sectional areas (strip thickness and bandwidth) can currently be produced, so that different alloy classes are necessary for different strength classes and / or cross-sectional areas.
  • the lowering of the carbon equivalent due to lower carbon and manganese contents is to be compensated for by increasing the silicon content.
  • the edge crack resistance and the weldability are improved with the same strength.
  • a low yield strength ratio (Re / Rm) in a strength range above 950 MPa in the initial state is typical for a dual-phase steel and is primarily used for the formability during stretching and deep-drawing processes. It gives the designer information about the distance between the onset of plastic deformation and failure of the material under quasi-static stress. Accordingly, lower yield strength ratios represent a greater safety margin from component failure.
  • a higher yield strength ratio (Re / Rm), as is typical for complex phase steels, is also characterized by a high resistance to edge cracks. This can be attributed to the smaller differences in the strength and hardness of the individual structural components and the finer structure, which has a favorable effect on a homogeneous deformation in the area of the cut edge.
  • the analytical landscape for achieving multi-phase steels with a minimum tensile strength of 950 MPa is very diverse and shows very large alloy ranges for the strength-increasing elements carbon, silicon, manganese, phosphorus, nitrogen, aluminum as well as chromium and / or molybdenum as well as in the addition of micro alloys such as Titanium, niobium, vanadium and boron.
  • the range of dimensions in this strength range is wide and lies in the thickness range from approximately 0.50 to approximately 4.00 mm for strips which are intended for continuous annealing.
  • Hot strip, cold-rolled hot strip and cold strip can be used as primary material. Tapes up to about 1600 mm wide are mainly used, but also Slit strip dimensions that result from slitting the strips lengthways. Sheets or sheets are made by cross-cutting the strips.
  • the structure of the steel is transferred to the austenitic area by heating, preferably to temperatures above 950 ° C. in a protective gas atmosphere. Subsequent cooling in air or protective gas leads to the formation of a martensitic structure for a high-strength component.
  • the subsequent tempering enables the reduction of residual stresses in the hardened component. At the same time, the hardness of the component is reduced so that the required toughness values are achieved.
  • the invention is therefore based on the object of creating a new cost-effective alloy concept for a high-strength air-hardenable multiphase steel with excellent processing properties and with a minimum tensile strength of 950 MPa in the non-tempered state, lengthways and crosswise to the rolling direction, preferably with a dual-phase structure, with which the process window for continuous annealing of hot or cold rolled strips has been expanded so that in addition to strips with different cross-sections, steel strips with a thickness and strip width that varies over the strip length and the correspondingly varying degrees of cold rolling with the most homogeneous mechanical and technological properties can be produced.
  • the hot-dip coating of the steel is to be guaranteed and a method for producing a strip made from this steel is to be specified.
  • the structure consists of the main phases ferrite and martensite and the secondary phase bainite, which determines the improved mechanical properties of the steel.
  • the steel is characterized by low carbon equivalents and, with the carbon equivalent CEV (IIW), is dependent on the sheet thickness for the addition of max. 0.66% limited, so that excellent weldability and the further specific properties described below can be achieved.
  • the steel Due to its chemical composition, the steel can be manufactured in a wide range of hot rolling parameters, for example with reel temperatures above the bainite start temperature (variant A).
  • a microstructure can be set which then allows the steel according to the invention to be cold rolled without prior soft annealing, with cold rolling degrees of between 10 and 40% being used per cold rolling pass.
  • the steel is very well suited as a primary material for hot-dip coating and, due to the sum-related amount of Mn, Si and Cr added according to the invention depending on the strip thickness to be produced, has a significantly enlarged process window compared to the known steels.
  • load-optimized components can advantageously be produced therefrom.
  • the steel strip according to the invention can be produced as cold and hot strip and as cold-rolled hot strip by means of a hot-dip galvanizing line or a pure continuous annealing system in the trained and undressed, in the stretch-bend-oriented and non-stretch-bend-oriented and also in the heat-treated (aged) state.
  • steel strips can be produced by an intercritical annealing between A c1 and A c3 or in the case of an austenitizing annealing over A c3 with a final controlled cooling, which leads to a dual or multi-phase structure.
  • Annealing temperatures of approximately 700 to 950 ° C. have proven to be advantageous. Depending on the overall process (only continuous annealing or additional hot-dip coating), there are different approaches to heat treatment.
  • the strip is cooled from the annealing temperature with a cooling rate of approx. 15 to 100 ° C / s to an intermediate temperature of approx. 160 to 250 ° C.
  • the cooling is stopped, as described above, before entering the molten bath and is continued only after the bath has exited until the intermediate temperature of about 200 to 250 ° C. has been reached.
  • the holding temperature in the molten bath is approximately 400 up to 470 ° C. Cooling down to room temperature takes place again at a cooling rate of approx. 2 to 30 ° C./s (see also method 2, Figure 6b ).
  • the second variant of the temperature control for hot-dip coating includes maintaining the temperature for approx. 1 to 20 s at the intermediate temperature of approx. 200 to 350 ° C and then reheating to the temperature required for hot-dip coating of approx. 400 to 470 ° C. After finishing, the strip is cooled again to approx. 200 to 250 ° C. The cooling to room temperature again takes place at a cooling rate of approx. 2 to 30 ° C./s (see also method 3, Figure 6c ).
  • manganese, chromium and silicon are responsible for the conversion of austenite to martensite in addition to carbon.
  • the carbon equivalent can be reduced, which improves the weldability and prevents excessive hardening during welding. In the case of resistance spot welding, the electrode service life can also be significantly increased.
  • Instruction elements are elements that are already present in the iron ore or, due to the manufacturing process, get into the steel. Because of their predominantly negative influences, they are usually undesirable. An attempt is made to remove them to a tolerable level or to convert them into harmless forms.
  • Hydrogen (H) is the only element that can diffuse through the iron lattice without generating lattice strain. This means that the hydrogen in the iron lattice is relatively mobile and can be absorbed relatively easily during the processing of the steel. Hydrogen can only be absorbed into the iron lattice in an atomic (ionic) form.
  • Hydrogen has a strong embrittlement effect and diffuses preferentially to energetically favorable places (defects, grain boundaries etc.). Defects act as hydrogen traps and can significantly increase the length of time that hydrogen remains in the material.
  • a more uniform structure which among other things in the steel according to the invention. achieved through its widened process window also reduces the susceptibility to hydrogen embrittlement.
  • Oxygen (O) In the molten state, the steel has a relatively high absorption capacity for gases. At room temperature, however, oxygen is only soluble in very small amounts. Analogous to hydrogen, oxygen can only diffuse into the material in an atomic form. Due to the strong embrittlement effect and the negative effects on the aging resistance, attempts are made to reduce the oxygen content as much as possible during manufacture.
  • the oxygen content in the steel should therefore be as low as possible.
  • Phosphorus (P) is a trace element from iron ore and is dissolved in the iron lattice as a substitute atom . Phosphorus increases hardness through solid-solution hardening and improves hardenability. However, attempts are generally made to lower the phosphorus content as much as possible, since, among other things, due to its low solubility in the solidifying medium, it tends to segregate and to a large extent reduces the toughness. Due to the accumulation of phosphorus at the grain boundaries, grain boundary breaks occur. In addition, phosphorus increases the transition temperature from tough to brittle behavior up to 300 ° C. During hot rolling, near-surface phosphorus oxides can cause tearing at the grain boundaries.
  • phosphorus is used in small quantities ( ⁇ 0.1% by weight) as a microalloying element due to the low cost and the high increase in strength, for example in high-strength IF steels (interstitial free), bake hardening steels or in some alloy concepts for dual phase steels.
  • the steel according to the invention differs from known analysis concepts which use phosphorus as a solid solution, inter alia in that phosphorus is not alloyed but is set as low as possible.
  • the phosphorus content in the steel according to the invention is limited to amounts which are unavoidable in the production of steel.
  • sulfur is bound as a trace element in iron ore.
  • Sulfur is undesirable in steel (with the exception of free-cutting steels) because it tends to segregate and has a strong embrittlement effect. An attempt is therefore made to achieve the lowest possible sulfur content in the melt, for example by means of a vacuum treatment.
  • the sulfur present is converted into the relatively harmless compound manganese sulfide (MnS) by adding manganese.
  • MnS manganese sulfide
  • the manganese sulfides are often rolled out in rows during the rolling process and act as germination points for the conversion. This leads to a stratified structure, especially in the case of diffusion-controlled conversion, and can lead to deteriorated mechanical properties in the case of pronounced stringency (e.g. pronounced marten seat lines instead of distributed martensite islands, anisotropic material behavior, reduced elongation at break).
  • the sulfur content in the steel according to the invention is limited to ⁇ 0.0030% by weight, advantageously to ⁇ 0.0025% by weight or optimally to ⁇ 0.0020% by weight or to quantities unavoidable in the production of steel .
  • Alloy elements are usually added to the steel in order to influence certain properties.
  • An alloy element in different steels can influence different properties. The effect generally depends strongly on the amount and the state of the solution in the material.
  • Carbon (C) is the most important alloying element in steel. Due to its targeted introduction of up to 2.06% by weight, iron only becomes steel. The carbon content is often drastically reduced during steel production. In the case of dual-phase steels for continuous hot-dip coating, its proportion according to EN 10346 or VDA 239-100 is a maximum of 0.230% by weight, a minimum value is not specified.
  • carbon is dissolved interstitially in the iron lattice.
  • the solubility is a maximum of 0.02% in ⁇ -iron and a maximum of 2.06% in ⁇ -iron.
  • carbon significantly increases the hardenability of steel and is therefore essential for the formation of a sufficient amount of martensite. Too high a carbon content, however, increases the difference in hardness between ferrite and martensite and limits weldability.
  • the steel according to the invention contains carbon contents of less than or equal to 0.115% by weight.
  • Silicon (Si) binds oxygen during casting and is therefore used for calming during the deoxidation of the steel. It is important for the later steel properties that the segregation coefficient is significantly lower than, for example, that of manganese (0.16 compared to 0.87). Segregations generally lead to a line arrangement of the structural components, which deteriorate the forming properties, for example the widening of the holes and the ability to bend.
  • the latter is due, among other things, to the fact that silicon reduces the solubility of carbon in the ferrite and increases the activity of carbon in the ferrite, thus preventing the formation of carbides, which, as brittle phases, reduce ductility, which in turn improves the formability.
  • the low strength-increasing effect of silicon within the range of the steel according to the invention creates the basis for a wide process window.
  • silicon in the range according to the invention has led to further surprising effects described below.
  • the delay in carbide formation described above could e.g. can also be brought about by aluminum.
  • aluminum forms stable nitrides, so that insufficient nitrogen is available for the formation of carbonitrides with microalloying elements.
  • This problem does not exist due to the alloying with silicon, since silicon forms neither carbides nor nitrides.
  • Silicon thus has an indirect positive effect on the formation of precipitates through microalloys, which in turn have a positive effect on the strength of the material. Since the increase in the transition temperatures due to silicon tends to favor grain coarsening, a microalloy with niobium, titanium and boron is particularly expedient, as is the targeted adjustment of the nitrogen content in the steel according to the invention.
  • the atmospheric conditions during the annealing treatment in a continuous hot-dip coating system result in a reduction in iron oxide, which is found, for example, in the Cold rolling or as a result of storage at room temperature on the surface.
  • the gas atmosphere is oxidizing, with the result that segregation and selective oxidation of these elements can occur.
  • the selective oxidation can take place both externally, that is to say on the substrate surface, and internally within the metallic matrix.
  • the internal oxidation of the alloying elements can be influenced in a targeted manner by adjusting the oxygen partial pressure of the furnace atmosphere (N 2 -H 2 protective gas atmosphere).
  • the set oxygen partial pressure must satisfy the following equation, the furnace temperature being between 700 and 950 ° C. - 12th > log pO 2nd ⁇ - 5 * Si - 0.25 - 3rd * Mn - 0.25 - 0.1 Cr - 0.5 - 7 * - ln B 0.5
  • Si, Mn, Cr, B denote the corresponding alloy proportions in the steel in% by weight and pO 2 the oxygen partial pressure in mbar.
  • the selective oxidation of the alloy elements can also be influenced via the gas atmospheres of the furnace areas.
  • the oxygen partial pressure and thus the oxidation potential for iron and the alloying elements can be set via the combustion reaction in the NOF. This must be set so that the oxidation of the alloy elements takes place internally below the steel surface and, if necessary, a thin iron oxide layer forms on the steel surface after the passage through the NOF area. This is achieved e.g. by reducing the CO value below 4% by volume.
  • the iron oxide layer that may be formed is reduced under an N2-H2 protective gas atmosphere and, likewise, the alloy elements are further oxidized internally.
  • the oxygen partial pressure set in this furnace area must satisfy the following equation, the furnace temperature being between 700 and 950 ° C. - 18th > log pO 2nd ⁇ - 5 * Si - 0.3 - 2.2 * Mn - 0.45 - 0.1 * Cr - 0.4 - 12.5 * - ln B 0.25
  • Si, Mn, Cr, B denote the corresponding alloy proportions in the steel in% by weight and pO 2 the oxygen partial pressure in mbar.
  • the dew point of the gas atmosphere N 2 -H 2 protective gas atmosphere
  • the oxygen partial pressure must be set so that oxidation of the strip before immersion in the molten bath is avoided. Dew points in the range of -30 to -40 ° C have proven to be advantageous.
  • hot-dip coating here, for example, hot-dip galvanizing
  • the process route is selected via continuous annealing with subsequent electrolytic galvanizing (see process 1 in Figure 6a )
  • electrolytic galvanizing pure zinc is deposited directly on the strip surface.
  • pure zinc is deposited directly on the strip surface.
  • it In order not to hinder the flow of electrons between the steel strip and the zinc ions and thus the galvanizing, it must be ensured that there is no surface-covering oxide layer on the strip surface. This condition is usually guaranteed by a standard reducing atmosphere during annealing and pre-cleaning before electrolysis.
  • the minimum silicon content is set at 0.400% by weight and the maximum silicon content at 0.500% by weight.
  • Manganese (Mn) is added to almost all steels for desulfurization in order to convert the harmful sulfur into manganese sulfides.
  • manganese increases the strength of the ferrite through solidification of the crystal and shifts the ⁇ - / ⁇ -conversion to lower temperatures.
  • the addition of manganese increases the hardness ratio between martensite and ferrite.
  • the structure of the structure is strengthened. A high difference in hardness between the phases and the formation of marten seat lines result in a lower hole expansion capacity, which is synonymous with increased sensitivity to edge cracking.
  • manganese tends to form oxides on the steel surface during the annealing treatment.
  • manganese oxides eg MnO
  • Mn mixed oxides eg Mn 2 SiO 4
  • Si / Mn or Al / Mn ratio manganese is to be regarded as less critical, since globular oxides form rather than oxide films.
  • high manganese levels can have a negative impact on the appearance of the zinc layer and the zinc adhesion.
  • the above-mentioned measures for setting the furnace areas during continuous hot dip coating reduce the formation of Mn oxides or Mn mixed oxides on the steel surface after annealing.
  • the manganese content is set at 1,900 to 2,350% by weight for the reasons mentioned.
  • the manganese content is preferably in a range between 1,9 1,900 and 2,2 2,200% by weight, with strip thicknesses of 1.00 to 2.00 mm between 2,0 2,050 and 50 2,250% by weight and for strip thicknesses over 2.00 mm between ⁇ 2,100% by weight and ⁇ 2,350% by weight.
  • Another special feature of the invention is that the variation in the manganese content can be compensated for by simultaneously changing the silicon content.
  • the coefficients of manganese and silicon are approximately the same for both the yield strength and the tensile strength, which makes it possible to replace manganese with silicon.
  • Chromium (Cr) on the one hand, can significantly increase the hardenability of steel in small quantities in dissolved form.
  • Cr Cr
  • chromium carbides causes particle solidification.
  • the associated increase in the number of germ sites with a simultaneously reduced carbon content leads to a reduction in the hardenability.
  • chromium In dual-phase steels, the addition of chromium mainly improves hardenability. When dissolved, chromium shifts the pearlite and bainite transformation for longer times and at the same time lowers the martensite start temperature.
  • Chromium is also a carbide former. If chromium-iron mixed carbides are present, the austenitizing temperature before hardening must be selected high enough to dissolve the chromium carbides. Otherwise, the increased number of bacteria can lead to a deterioration in the hardenability.
  • Chromium also tends to form oxides on the steel surface during the annealing treatment, which can degrade the hot dip quality.
  • the above-mentioned measures for setting the furnace areas during continuous hot dip coating reduce the formation of Cr oxides or Cr mixed oxides on the steel surface after annealing.
  • the chromium content is therefore set at contents of 0.200 to 0.500% by weight.
  • Molybdenum (Mo) The addition of molybdenum leads to an improvement in hardenability, similar to that of chromium and manganese. The pearlite and bainite transformation is shifted to longer times and the martensite start temperature is lowered. At the same time, molybdenum is a strong chalk former, which produces finely divided mixed carbides, including with titanium. Molybdenum also significantly increases the tempering resistance, so that no loss of strength is to be expected in the hot-dip bath. Molybdenum also works through mixed crystal hardening, but is less effective than manganese and silicon.
  • the molybdenum content is therefore set between 0.200 to 0.300% by weight. Ranges between 0.200 and 0.250% by weight are advantageous.
  • Copper (Cu) The addition of copper can increase tensile strength and hardenability. In combination with nickel, chromium and phosphorus, copper can form a protective oxide layer on the surface, which can significantly reduce the rate of corrosion.
  • copper In combination with oxygen, copper can form harmful oxides at the grain boundaries, which can have negative effects especially for hot forming processes.
  • the copper content is therefore set at ⁇ 0.050% by weight and is therefore limited to the amounts that are unavoidable in steel production.
  • the nickel content is therefore set at ⁇ 0.050% by weight and is therefore limited to the amounts that are unavoidable in steel production.
  • Vanadium (V) Since the addition of vanadium is not necessary in the present alloy concept, the vanadium content is limited to inevitable amounts accompanying the steel.
  • Aluminum (Al) is usually alloyed to the steel in order to bind the oxygen and nitrogen dissolved in the iron. Oxygen and nitrogen are thus converted into aluminum oxides and aluminum nitrides. These excretions can cause grain refinement by increasing the number of germs and thus increase the toughness properties and strength values.
  • Titanium nitrides have a lower enthalpy of formation and are formed at higher temperatures.
  • the aluminum content is therefore limited to 0.005 to a maximum of 0.060% by weight and is added to calm the steel.
  • Niobium acts in steel in different ways. When hot rolling in the finishing train, it delays recrystallization due to the formation of very finely divided precipitates, which increases the density of germination points and results in a finer grain after conversion. The proportion of dissolved niobium also inhibits recrystallization. The excretions increase strength in the final product. These can be carbides or carbonitrides. Often it is mixed carbides, in which titanium is also incorporated. This effect starts from 0.005% by weight and is most evident from 0.010% by weight of niobium. The precipitates also prevent grain growth during (partial) austenitization in hot-dip galvanizing. No additional effect is to be expected above 0.060% by weight of niobium. Contents of 0.025 to 0.045% by weight have proven to be advantageous.
  • Titanium (Ti) Due to its high affinity for nitrogen, titanium is primarily excreted as TiN during solidification. It also occurs together with niobium as a mixed carbide. TiN is of great importance for grain size stability in the pusher furnace. The Excretions have a high temperature stability, so that, in contrast to the mixed carbides at 1200 ° C, they are mostly present as particles that hinder grain growth. Titanium also retards recrystallization during hot rolling, but is less effective than niobium. Titan works through precipitation hardening. The larger TiN particles are less effective than the more finely distributed mixed carbides. The best effectiveness is achieved in the range from 0.005 to 0.060% by weight of titanium, which is why this represents the alloy range according to the invention. For this, contents of 0.025 to 0.045% by weight have been found to be advantageous.
  • Boron (B) Boron is an extremely effective alloying agent to increase hardenability, which is effective even in very small amounts (from 5 ppm). The martensite start temperature remains unaffected.
  • boron must be in solid solution. Since it has a high affinity for nitrogen, the nitrogen must first be set, preferably by the stoichiometrically necessary amount of titanium. Due to its low solubility in iron, the dissolved boron preferentially attaches to the austenite grain boundaries. There it partially forms Fe-B carbides, which are coherent and reduce the grain boundary energy. Both effects delay the formation of ferrite and pearlite and thus increase the hardenability of the steel.
  • the boron content for the alloy concept according to the invention is set at values from 5 to 30 ppm, advantageously at ⁇ 25 or optimally at ⁇ 20 ppm.
  • Nitrogen (N) can be an alloying element as well as an accompanying element from steel production. Too high levels of nitrogen lead to an increase in strength combined with a rapid loss of toughness and aging effects.
  • fine grain hardening can be achieved using titanium nitride and niobium (carbo) nitride. Coarse grain formation is also suppressed when reheating before hot rolling.
  • the N content is therefore set to values of 0,00 0.0020 to ⁇ 0.0120% by weight.
  • the nitrogen content should be maintained at values of ⁇ 20 to ⁇ 90 ppm.
  • nitrogen contents of ⁇ 40 to ⁇ 120 ppm have proven to be advantageous.
  • niobium and titanium contents of ⁇ 0.100% by weight have proven to be advantageous and because of the principle interchangeability of niobium and titanium up to a minimum niobium content of 10 ppm and, for reasons of cost, particularly advantageous of ⁇ 0.090% by weight.
  • total contents of 0 0.102% by weight have proven to be advantageous and particularly advantageous ⁇ 0.092% by weight. Higher contents no longer have an improvement in the sense of the invention.
  • Calcium (Ca) An addition of calcium in the form of calcium-silicon mixed compounds causes a deoxidation and desulfurization of the molten phase during the production of steel. In this way, reaction products are transferred to the slag and the steel is cleaned. The increased purity leads to better properties according to the invention in the end product.
  • the annealing temperatures for the dual-phase structure to be achieved for the steel according to the invention are between approximately 700 and 950 ° C., so that depending on the temperature range, a partially austenitic (two-phase area) or a fully austenitic structure (austenite area) is achieved.
  • the continuously annealed and, in some cases, hot-dip coated material can be manufactured both as hot strip and as cold-rolled hot strip or cold strip in the trained (cold-rolled) or undressed state and / or in the stretch-oriented or non-stretch-bent state and also in the heat-treated state (aging). This state is referred to below as the initial state.
  • Steel strips in the present case as hot strip, cold-rolled hot strip or cold strip, from the alloy composition according to the invention are also distinguished by a high resistance to edge cracking during further processing.
  • the hot strip is produced according to the invention with finish rolling temperatures in the austenitic region above A r3 and at reel temperatures above the bainite start temperature (variant A).
  • the hot strip is produced according to the invention with finish rolling temperatures in the austenitic region above A r3 and coiling temperatures below the bainite start temperature (variant B).
  • Figure 1 shows schematically the process chain for the production of a strip from the steel according to the invention.
  • the different process routes relating to the invention are shown.
  • the process route is the same for all steels according to the invention until hot rolling (final rolling temperature), after which process routes differ depending on the desired results.
  • the pickled hot strip can be galvanized or cold rolled and galvanized with different degrees of rolling.
  • Soft-annealed hot strip or soft-annealed cold strip can also be cold-rolled and galvanized.
  • material can also be processed without hot-dip coating, i.e. only in the context of continuous annealing with and without subsequent electrolytic galvanizing.
  • a complex component can now be produced from the optionally coated material. This is followed by the hardening process, in which the air is cooled in accordance with the invention.
  • a tempering stage can complete the thermal treatment of the component.
  • Figure 2 shows schematically the time-temperature profile of the process steps hot rolling and continuous annealing of strips from the alloy composition according to the invention. It shows the time and temperature-dependent conversion for the hot rolling process as well as for heat treatment after cold rolling, component production, tempering and optional tempering.
  • Figure 3 shows the chemical composition of the investigated steels in the upper half of the table. Alloys LH®1100 according to the invention were compared with the reference grades LH®800 / LH®900.
  • the alloys according to the invention in particular have significantly higher Si contents and lower Cr contents and no V alloy.
  • Figure 4 shows the mechanical parameters along the rolling direction of the investigated steels, with target values to be achieved for the air-hardened state ( Figure 4a ), the determined Values in the non-air-hardened initial state ( Figure 4b ) and in air-hardened condition ( Figure 4c ). The specified values to be achieved are safely achieved.
  • Figure 5 shows results of hole expansion tests according to ISO 16630 (absolute values).
  • the results of the hole expansion tests for variant A are shown for method 2 ( Figure 6b , 1 , 2nd mm) and method 3 ( Figure 6c , 2.0 mm).
  • the investigated materials have a sheet thickness of 1.2 or 2.0 mm.
  • the results apply to the test according to ISO 16630.
  • Method 2 corresponds to annealing, for example on hot-dip galvanizing with a combined direct-fired furnace and radiant tube furnace, as described in Figure 6b is described.
  • the method 3 corresponds, for example, to a process control in a continuous annealing system as shown in Figure 6c is described.
  • the steel can optionally be reheated directly in front of the zinc bath using an induction furnace.
  • the Figure 6 shows schematically three variants of the temperature-time profiles according to the invention in the annealing treatment and cooling and in each case different austenitizing conditions.
  • Procedure 1 shows the annealing and cooling of the cold or hot-rolled or cold-rolled steel strip produced in a continuous annealing plant.
  • the tape is heated to a temperature in the range of about 700 to 950 ° C (Ac1 to Ac3).
  • the annealed steel strip is then cooled from the annealing temperature with a cooling rate between approx. 15 and 100 ° C / s to an intermediate temperature (ZT) of approx. 200 to 250 ° C.
  • ZT intermediate temperature
  • This schematic representation does not show a second intermediate temperature (approx. 300 to 500 ° C).
  • the steel strip is cooled at a cooling rate of between about 2 and 30 ° C / s until reaching the R aum t emperature (RT) in air or cooling at a cooling rate between about 15 and 100 ° C / s up maintain at room temperature.
  • RT R aum t emperature
  • Procedure 2 shows the process according to method 1, however, the cooling of the steel strip is temporarily interrupted for the purpose of hot-dip coating when it passes through the hot-dip tank, in order to then cool at a cooling rate between approx. 15 and 100 ° C / s up to an intermediate temperature of approx. 200 continue up to 250 ° C.
  • the steel strip is then cooled in air at a cooling rate between approx. 2 and 30 ° C / s until room temperature is reached.
  • Procedure 3 ( Figure 6c ) also shows the process according to method 1 for hot-dip coating, but the cooling of the steel strip is interrupted by a short pause (approx. 1 to 20 s) at an intermediate temperature in the range of approx. 200 to 400 ° C and up to the temperature ( ST), which is necessary for hot-dip coating (approx. 400 to 470 ° C), reheated.
  • the steel strip is then cooled again to an intermediate temperature of approx. 200 to 250 ° C.
  • the final cooling of the steel strip takes place at a cooling rate of approx. 2 and 30 ° C / s until the room temperature is reached in air.
  • Example 1 (cold strip) (alloy composition in% by weight)
  • the material was previously hot-rolled at a final rolling set temperature of 910 ° C and coiled at a reel set temperature of 650 ° C with a thickness of 2.30 mm and after Pickling without additional heat treatment (such as hood annealing) cold rolled twice with an intermediate thickness of 1.49 mm.
  • the steel according to the invention After tempering, the steel according to the invention has a structure which consists of martensite, bainite and residual austenite.
  • This steel shows the following characteristic values after air hardening (initial values in brackets, unrefined condition): - yield strength (Rp0.2) 921 MPa (768 MPa) - tensile strength (Rm) 1198 MPa (984 MPa) - elongation at break (A80) 5.5% (10.7%) - A5 stretch 9.5% (-) - Bake hardening index (BH2) 52 MPa - Hole expansion ratio according to ISO 16630 - (49%) - Bending angle according to VDA 238-100 (lengthways, crossways) - (122 ° / 112 °) longitudinal to the rolling direction and would correspond to an LH®1100, for example.
  • the yield point ratio Re / Rm in the longitudinal direction was 78% in the initial state.
  • Example 2 (cold strip) (alloy composition in% by weight)
  • This steel shows the following characteristic values after air hardening (initial values in brackets, unrefined condition): - yield strength (Rp0.2) 903 MPa (708 MPa) - tensile strength (Rm) 1186 MPa (983 MPa) - elongation at break (A80) 7.1% (11.7%) - A5 stretch 9.1% (-) - Bake hardening index (BH2) 48 MPa - Hole expansion ratio according to ISO 16630 - (32%) - Bending angle according to VDA 238-100 (lengthways, crossways) - (104 ° / 88 °) longitudinal to the rolling direction and would correspond to an LH®1100, for example. The yield point ratio Re / Rm in the longitudinal direction was 72% in the initial state.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Chemical Kinetics & Catalysis (AREA)
  • Oil, Petroleum & Natural Gas (AREA)
  • Heat Treatment Of Sheet Steel (AREA)

Claims (30)

  1. Bande d'acier laminée à froid ou à chaud en un acier à plusieurs phases pouvant être trempé à l'air et revenu, présentant une résistance minimale à la traction dans le sens longitudinal et le sens transversal par rapport à la direction du laminage de 950 MPa avant la trempe à l'air et le revenu, présentant d'excellentes propriétés de mise en œuvre, consistant en les éléments suivants (teneurs en % en poids) :
    C ≥ 0,075 à ≤ 0,115
    Si ≥ 0,400 à ≤ 0,500
    Mn ≥ 1,900 à ≤ 2,350
    Cr ≥ 0,200 à ≤ 0,500
    Al ≥ 0,005 à ≤ 0,060
    N ≥ 0,0020 à ≤ 0,0120
    S ≤ 0,0030
    Nb ≥ 0,005 à ≤ 0,060
    Ti ≥ 0,005 à ≤ 0,060
    B ≥ 0,0005 à ≤ 0,0030
    Mo ≥ 0,200 à ≤ 0,300
    Ca ≥ 0,0005 à ≤ 0,0060
    Cu ≤ 0,050
    Ni ≤ 0,050
    le reste étant constitué de fer, y compris les impuretés usuelles d'accompagnement dans l'acier dues à l'élaboration, dans laquelle, eu égard à une fenêtre de procédé aussi large que possible lors du recuit continu de bandes à chaud ou à froid en cet acier, la teneur totale en Mn+Si+Cr est ajustée comme suit en fonction de l'épaisseur de la bande à produire :
    jusqu'à 1,00 mm : somme de Mn+Si+Cr ≥ 2,800 et ≤ 3,000 % en poids,
    supérieure à 1,00 et allant jusqu'à 2,00 mm : somme de Mn+Si+Cr ≥ 2,850 et ≤ 3,100 % en poids,
    supérieure à 2,00 mm : somme de Mn+Si+Cr ≥ 2,900 et ≤ 3,200 % en poids.
  2. Bande d'acier selon la revendication 1, caractérisée en ce que, pour des épaisseurs de bande allant jusqu'à 1,00 mm, la teneur en C est ≤ 0,100 % et l'équivalent carbone CEV(IIW) est ≤ 0,62 % en poids.
  3. Bande d'acier selon la revendication 1, caractérisée en ce que, pour des épaisseurs de bande supérieures à 1,00 et allant jusqu'à 2,00 mm, la teneur en C est ≤ 0,105 % et l'équivalent carbone CEV(IIW) est ≤ 0,64 %.
  4. Bande d'acier selon la revendication 1, caractérisée en ce que, pour des épaisseurs de bande supérieures à 2,00 mm, la teneur en C est ≤ 0,115 % et l'équivalent carbone CEV(IIW) est ≤ 0,66 %.
  5. Bande d'acier selon les revendications 1 et 2, caractérisée en ce que, pour des épaisseurs de bande allant jusqu'à 1,00 mm, la teneur en Mn est ≥ 1,900 jusqu'à ≤ 2,200 %.
  6. Bande d'acier selon les revendications 1 et 3, caractérisée en ce que, pour des épaisseurs de bande supérieures à 1,00 et allant jusqu'à 2,00 mm, la teneur en Mn est ≥ 2,050 jusqu'à ≤ 2,250 %.
  7. Bande d'acier selon les revendications 1 et 4, caractérisée en ce que, pour des épaisseurs de bande supérieures à 2,00 mm, la teneur en Mn est ≥ 2,100 jusqu'à ≤ 2,350 %.
  8. Bande d'acier selon l'une des revendications 1 à 7, caractérisée en ce que, pour une somme de Ti+Nb+B ≥ 0,010 jusqu'à ≤ 0,070 %, la teneur en N est ≥ 0,0020 jusqu'à ≤ 0,0090 %.
  9. Bande d'acier selon la revendication 8, caractérisée en ce que, pour la somme de Ti+Nb+B > 0,070 %, la teneur en N est ≥ 0,0040 jusqu'à ≤ 0,0120 %.
  10. Bande d'acier selon l'une des revendications 1 à 9, caractérisée en ce que la teneur en S est ≤ 0,0025 %.
  11. Bande d'acier selon la revendication 10, caractérisée en ce que la teneur en S est ≤ 0,0020 %.
  12. Bande d'acier selon l'une des revendications 1 à 11, caractérisée en ce que la teneur en Mo est ≤ 0,250 %.
  13. Bande d'acier selon l'une des revendications 1 à 12, caractérisée en ce que la teneur en Ti est ≥ 0,025 ≤ 0,045 %.
  14. Bande d'acier selon l'une des revendications 1 à 13, caractérisée en ce que la teneur en Nb est ≥ 0,025 jusqu'à ≤ 0,045 %.
  15. Bande d'acier selon l'une des revendications 1 à 14, caractérisée en ce que la somme Nb+Ti ≤ 0,100 %.
  16. Bande d'acier selon la revendication 15, caractérisée en ce que la somme Nb+Ti est ≤ 0,090 %.
  17. Bande d'acier selon l'une des revendications 1 à 16, caractérisée en ce que la somme Cr+Mo est ≤ 0,725 %.
  18. Bande d'acier selon l'une des revendications 1 à 17, caractérisée en ce que la somme Ti+Nb+B est ≤ 0,102 %.
  19. Bande d'acier selon la revendication 18, caractérisée en ce que la somme Ti+Nb+B est ≤ 0,092 %.
  20. Bande d'acier selon l'une des revendications 1 à 19, caractérisée en ce que la teneur en Ca est ≤ 0,0030 %.
  21. Traitement thermique d'une bande d'acier laminée à froid ou à chaud en un acier à plusieurs phases pouvant subir une trempe à l'air et un recuit selon l'une des revendications 1 à 20, caractérisé en ce que la bande d'acier laminée à chaud ou à froid est, pendant le recuit continu, chauffée à une température comprise dans la plage d'environ 700 à 950 °C, et que la bande d'acier recuite est ensuite refroidie, à partir de la température de recuit, à une vitesse de refroidissement comprise entre environ 15 et 100 °C/s jusqu'à une première température intermédiaire d'environ 300 à 500 °C, puis à une vitesse de refroidissement comprise entre environ 15 et 100 °C/s jusqu'à une seconde température intermédiaire d'environ 160 à 250 °C, puis la bande d'acier est refroidie à l'air à une température de refroidissement d'environ 2 à 30 °C/s jusqu'à atteindre la température ambiante, ou le refroidissement est maintenu à une vitesse de refroidissement comprise entre environ 15 et 100 °C/s, de la première température intermédiaire à la température ambiante.
  22. Traitement thermique selon la revendication 21, caractérisé en ce que, lors d'un revêtement par immersion à chaud, après le chauffage, suivi d'un refroidissement, le refroidissement est interrompu avant pénétration dans le bain de fusion et, après le revêtement par immersion à chaud, on poursuit le refroidissement à une vitesse de refroidissement comprise entre environ 15 et 100 °C/s jusqu'à une température intermédiaire d'environ 200 à 250 °C, puis la bande d'acier est refroidie à l'air à une vitesse de refroidissement d'environ 2 à 30 °C/s jusqu'à atteindre la température ambiante.
  23. Traitement thermique selon la revendication 1, caractérisé en ce que, lors d'un revêtement par immersion à chaud, après le chauffage, puis le refroidissement à la température intermédiaire d'environ 200 à 250 °C, on maintient la température pendant environ 1 à 20 s avant pénétration dans le bain de fusion, puis on chauffe de nouveau la bande d'acier à une température d'environ 400 à 470 °C et, après achèvement du revêtement par immersion à chaud, on procède à un refroidissement à une température de refroidissement comprise entre environ 15 et 100 °C/s jusqu'à une température intermédiaire d'environ 200 à 250 °C, puis on refroidit à l'air jusqu'à la température ambiante à une vitesse de refroidissement d'environ 2 à 30 °C/s.
  24. Traitement thermique selon l'une des revendications 22 à 23, caractérisé en ce que, lors du recuit continu, on élève le potentiel d'oxydation lors d'un recuit avec une configuration d'installation consistant en une zone de four à flamme directe (NOF) et un four à tubes radiants (RTF) grâce à une teneur en CO dans le NOF d'environ 4 % en volume, la pression partielle d'oxygène de l'atmosphère du four réductrice pour le fer étant, dans le RTF, ajustée selon l'équation suivante : − 18 > Log pO 2 ≥ − 5 * Si − 0,3 − 2,2 * Mn − 0,45 − 0,1 * Cr − 0,4 − 12,5 * ln B 0,25
    Figure imgb0010
     Si, Mn, Cr, B désignant les proportions d'alliage correspondantes, dans l'acier en % en poids, et p02 désignant la pression partielle d'oxygène en mbar et, pour éviter l'oxydation de la bande immédiatement avant immersion dans le bain de fusion, on ajuste le point de rosée de l'atmosphère gazeuse à -30 °C ou moins.
  25. Traitement thermique selon l'une des revendications 22 à 23, caractérisé en ce que, lors d'un recuit, seulement dans un four à tubes radiants, la pression partielle d'oxygène de l'atmosphère du four satisfait à l'équation suivante : − 12 > Log pO 2 ≥ − 5 * Si − 0,25 − 3 * Mn − 0,5 − 0,1 * Cr − 0,5 − 7 * ln B 0,5
    Figure imgb0011
     Si, Mn, Cr, B désignant les proportions d'alliage correspondantes, dans l'acier en % en poids, et p02 désignant la pression partielle d'oxygène en mbar et, pour éviter l'oxydation de la bande immédiatement avant immersion dans le bain de fusion, on ajuste le point de rosée de l'atmosphère gazeuse à -30 °C ou moins.
  26. Procédé selon l'une des revendications 21 à 25, caractérisé en ce que la bande d'acier est dressée après le traitement thermique ou le revêtement par immersion à chaud.
  27. Procédé selon au moins l'une des revendications 21 à 26, caractérisé en ce que la bande d'acier est dressée par traction et flexion après le traitement thermique ou le revêtement par immersion à chaud.
  28. Bande d'acier fabriquée par le procédé selon au moins l'une des revendications 21 à 27, présentant un rapport minimal d'expansion de trou selon ISO 16630 de 25 %.
  29. Bande d'acier fabriquée par le procédé selon au moins l'une des revendications 21 à 27, présentant un angle minimal de pliage selon VDA 238-100 de 65° dans la direction longitudinale ou dans la direction transversale.
  30. Bande d'acier fabriquée par le procédé selon au moins l'une des revendications 21 à 27, présentant un produit minimal Rm x α (résistance à la traction x angle de pliage selon VDA 238-100) de 100 000 MPa.
EP15821018.7A 2014-11-18 2015-11-06 Acier multiphases autotrempant en profondeur à haute résistance mécanique avec excellent aptitude au faconnage et procédé de fabrication d'une bande de cet acier Active EP3221483B1 (fr)

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
DE102014017274.0A DE102014017274A1 (de) 2014-11-18 2014-11-18 Höchstfester lufthärtender Mehrphasenstahl mit hervorragenden Verarbeitungseigenschaften und Verfahren zur Herstellung eines Bandes aus diesem Stahl
PCT/DE2015/100474 WO2016078644A1 (fr) 2014-11-18 2015-11-06 Acier polyphasé, trempé à l'air et à haute résistance, ayant d'excellentes propriétés de mise en oeuvre et procédé de production d'une bande avec cet acier

Publications (2)

Publication Number Publication Date
EP3221483A1 EP3221483A1 (fr) 2017-09-27
EP3221483B1 true EP3221483B1 (fr) 2020-05-06

Family

ID=55077320

Family Applications (1)

Application Number Title Priority Date Filing Date
EP15821018.7A Active EP3221483B1 (fr) 2014-11-18 2015-11-06 Acier multiphases autotrempant en profondeur à haute résistance mécanique avec excellent aptitude au faconnage et procédé de fabrication d'une bande de cet acier

Country Status (8)

Country Link
US (1) US10626478B2 (fr)
EP (1) EP3221483B1 (fr)
KR (1) KR20170084210A (fr)
CN (1) CN107208232B (fr)
DE (1) DE102014017274A1 (fr)
MX (1) MX2017006374A (fr)
RU (1) RU2721767C2 (fr)
WO (1) WO2016078644A1 (fr)

Cited By (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP3825433B1 (fr) * 2018-08-22 2023-02-15 JFE Steel Corporation Tôle d'acier de haute résistance et méthode de production pour celle-ci
EP3950994B1 (fr) * 2019-03-28 2024-01-24 Nippon Steel Corporation Tôle d'acier à haute résistance

Families Citing this family (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
DE102015111177A1 (de) * 2015-07-10 2017-01-12 Salzgitter Flachstahl Gmbh Höchstfester Mehrphasenstahl und Verfahren zur Herstellung eines kaltgewalzten Stahlbandes hieraus
EP3571324B1 (fr) * 2017-01-20 2021-11-03 thyssenkrupp Hohenlimburg GmbH Produit plat en acier laminé à chaud constitué d'un acier à phase complexe ayant une structure essentiellement bainitique et procédé pour fabriquer un tel produit plat en acier
DE102017123236A1 (de) 2017-10-06 2019-04-11 Salzgitter Flachstahl Gmbh Höchstfester Mehrphasenstahl und Verfahren zur Herstellung eines Stahlbandes aus diesem Mehrphasenstahl
DE102017131253A1 (de) * 2017-12-22 2019-06-27 Voestalpine Stahl Gmbh Verfahren zum Erzeugen metallischer Bauteile mit angepassten Bauteileigenschaften
KR102507715B1 (ko) 2018-08-22 2023-03-07 제이에프이 스틸 가부시키가이샤 고강도 강판 및 그의 제조 방법
US11795518B2 (en) * 2018-10-24 2023-10-24 Nippon Steel Corporation Non-oriented electrical steel sheet and method of manufacturing stacked core using same
DE102020110319A1 (de) 2020-04-15 2021-10-21 Salzgitter Flachstahl Gmbh Verfahren zur Herstellung eines Stahlbandes mit einem Mehrphasengefüge und Stahlband hinzu

Citations (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
DE102012013113A1 (de) * 2012-06-22 2013-12-24 Salzgitter Flachstahl Gmbh Hochfester Mehrphasenstahl und Verfahren zur Herstellung eines Bandes aus diesem Stahl mit einer Mindestzugfestigkleit von 580MPa

Family Cites Families (58)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4578124A (en) * 1984-01-20 1986-03-25 Kabushiki Kaisha Kobe Seiko Sho High strength low carbon steels, steel articles thereof and method for manufacturing the steels
JP3347151B2 (ja) * 1991-11-18 2002-11-20 日新製鋼株式会社 耐食性に優れた低降伏比冷延高張力鋼板の製造方法
EP0565066B1 (fr) * 1992-04-06 1997-07-02 Kawasaki Steel Corporation TÔle noire ou fer blanc pour la production de boîtes et procédé de fabrication
DE19610675C1 (de) 1996-03-19 1997-02-13 Thyssen Stahl Ag Mehrphasenstahl und Verfahren zu seiner Herstellung
DE10037867A1 (de) 1999-08-06 2001-06-07 Muhr & Bender Kg Verfahren zum flexiblen Walzen eines Metallbandes
WO2001064967A1 (fr) * 2000-02-29 2001-09-07 Kawasaki Steel Corporation Tole d'acier laminee a froid a haute resistance presentant d'excellentes proprietes de durcissement par vieillissement par l'ecrouissage
CA2422753C (fr) * 2000-09-21 2007-11-27 Nippon Steel Corporation Tole d'acier presentant de bonnes caracteristiques de gel de forme et procede permettant de produire cette tole
JP2002173742A (ja) * 2000-12-04 2002-06-21 Nisshin Steel Co Ltd 形状平坦度に優れた高強度オーステナイト系ステンレス鋼帯およびその製造方法
FR2847273B1 (fr) * 2002-11-19 2005-08-19 Usinor Piece d'acier de construction soudable et procede de fabrication
JP4005517B2 (ja) * 2003-02-06 2007-11-07 株式会社神戸製鋼所 伸び、及び伸びフランジ性に優れた高強度複合組織鋼板
CN100371487C (zh) * 2003-04-28 2008-02-27 杰富意钢铁株式会社 盘式制动器用马氏体类不锈钢
AU2003235443A1 (en) * 2003-05-27 2005-01-21 Nippon Steel Corporation High strength thin steel sheet excellent in resistance to delayed fracture after forming and method for preparation thereof, and automobile parts requiring strength manufactured from high strength thin steel sheet
JP4443910B2 (ja) * 2003-12-12 2010-03-31 Ｊｆｅスチール株式会社 自動車構造部材用鋼材およびその製造方法
DE102004053620A1 (de) 2004-11-03 2006-05-04 Salzgitter Flachstahl Gmbh Hochfester, lufthärtender Stahl mit ausgezeichneten Umformeigenschaften
US7442268B2 (en) * 2004-11-24 2008-10-28 Nucor Corporation Method of manufacturing cold rolled dual-phase steel sheet
EP1867748A1 (fr) * 2006-06-16 2007-12-19 Industeel Creusot Acier inoxydable duplex
CN101426942A (zh) * 2006-10-05 2009-05-06 杰富意钢铁株式会社 抗回火软化性和韧性优良的制动盘
KR100851189B1 (ko) * 2006-11-02 2008-08-08 주식회사 포스코 저온인성이 우수한 초고강도 라인파이프용 강판 및 그제조방법
EP1990431A1 (fr) * 2007-05-11 2008-11-12 ArcelorMittal France Procédé de fabrication de tôles d'acier laminées à froid et recuites à très haute résistance, et tôles ainsi produites
PL2031081T3 (pl) 2007-08-15 2011-11-30 Thyssenkrupp Steel Europe Ag Stal dwufazowa, produkt płaski z takiej stali dwufazowej i sposób wytwarzania produktu płaskiego
EP2028282B1 (fr) 2007-08-15 2012-06-13 ThyssenKrupp Steel Europe AG Acier en phase double, produit plat à partir d'un tel acier en phase double et son procédé de fabrication
DE102007058222A1 (de) 2007-12-03 2009-06-04 Salzgitter Flachstahl Gmbh Stahl für hochfeste Bauteile aus Bändern, Blechen oder Rohren mit ausgezeichneter Umformbarkeit und besonderer Eignung für Hochtemperatur-Beschichtungsverfahren
JP4894863B2 (ja) * 2008-02-08 2012-03-14 Ｊｆｅスチール株式会社 加工性に優れた高強度溶融亜鉛めっき鋼板およびその製造方法
EP2309014B1 (fr) * 2008-07-31 2013-12-25 JFE Steel Corporation Tôles d'acier épaisses laminées à chaud présentant une résistance élevée à la traction et une excellente résistance à basse température, et procédé de production de celles-ci
EP2163659B1 (fr) * 2008-09-11 2016-06-08 Outokumpu Nirosta GmbH Acier inoxidable, bande froide fabriquée à partir de cet acier et procédé de fabrication d'un produit plat en acier à partir de cet acier
JP5438302B2 (ja) * 2008-10-30 2014-03-12 株式会社神戸製鋼所 加工性に優れた高降伏比高強度の溶融亜鉛めっき鋼板または合金化溶融亜鉛めっき鋼板とその製造方法
CA2809171C (fr) * 2009-01-30 2017-12-19 Jfe Steel Corporation Tole forte d'acier laminee a chaud a resistance elevee a la traction presentant une excellente resistance aux fissurations induites par l'hydrogene
JP4924730B2 (ja) * 2009-04-28 2012-04-25 Ｊｆｅスチール株式会社 加工性、溶接性および疲労特性に優れる高強度溶融亜鉛めっき鋼板およびその製造方法
DE102009030489A1 (de) * 2009-06-24 2010-12-30 Thyssenkrupp Nirosta Gmbh Verfahren zum Herstellen eines warmpressgehärteten Bauteils, Verwendung eines Stahlprodukts für die Herstellung eines warmpressgehärteten Bauteils und warmpressgehärtetes Bauteil
DE102010024664A1 (de) 2009-06-29 2011-02-17 Salzgitter Flachstahl Gmbh Verfahren zum Herstellen eines Bauteils aus einem lufthärtbaren Stahl und ein damit hergestelltes Bauteil
CN102712171B (zh) * 2009-12-21 2015-03-25 塔塔钢铁艾默伊登有限责任公司 高强度的热浸镀锌钢带材
EP2524972B9 (fr) * 2010-01-13 2017-08-30 Nippon Steel & Sumitomo Metal Corporation Procede de production d'une tole d'acier a haute resistance presentant une excellente formabilite.
JP5434960B2 (ja) * 2010-05-31 2014-03-05 Ｊｆｅスチール株式会社 曲げ性および溶接性に優れる高強度溶融亜鉛めっき鋼板およびその製造方法
ES2455222T5 (es) * 2010-07-02 2018-03-05 Thyssenkrupp Steel Europe Ag Acero de resistencia superior, conformable en frío y producto plano de acero compuesto de un acero de este tipo
JP5306418B2 (ja) * 2010-07-09 2013-10-02 日新製鋼株式会社 銅被覆鋼箔、負極用電極及び電池
US9512499B2 (en) * 2010-10-22 2016-12-06 Nippon Steel & Sumitomo Metal Corporation Method for manufacturing hot stamped body having vertical wall and hot stamped body having vertical wall
KR101533164B1 (ko) * 2010-10-22 2015-07-01 신닛테츠스미킨 카부시키카이샤 핫스탬프 성형체의 제조 방법 및 핫스탬프 성형체
US9896736B2 (en) * 2010-10-22 2018-02-20 Nippon Steel & Sumitomo Metal Corporation Method for manufacturing hot stamped body having vertical wall and hot stamped body having vertical wall
DE102011000089A1 (de) * 2011-01-11 2012-07-12 Thyssenkrupp Steel Europe Ag Verfahren zum Herstellen eines warmgewalzten Stahlflachprodukts
DE102011117572A1 (de) * 2011-01-26 2012-08-16 Salzgitter Flachstahl Gmbh Höherfester Mehrphasenstahl mit ausgezeichneten Umformeigenschaften
US9670569B2 (en) * 2011-03-28 2017-06-06 Nippon Steel & Sumitomo Metal Corporation Cold-rolled steel sheet and production method thereof
EP2524970A1 (fr) * 2011-05-18 2012-11-21 ThyssenKrupp Steel Europe AG Produit plat en acier hautement résistant et son procédé de fabrication
MX2014002922A (es) * 2011-09-13 2014-05-21 Tata Steel Ijmuiden Bv Tira de acero galvanizada por inmersion en caliente de alta resistencia.
CN103987868B (zh) * 2011-09-30 2016-03-09 新日铁住金株式会社 具有980MPa以上的最大拉伸强度、材质各向异性少且成形性优异的高强度热浸镀锌钢板、高强度合金化热浸镀锌钢板及它们的制造方法
EP2799562B1 (fr) * 2011-12-27 2015-10-14 JFE Steel Corporation Feuille d'acier laminée à chaud et son procédé de fabrication
DE102012002079B4 (de) 2012-01-30 2015-05-13 Salzgitter Flachstahl Gmbh Verfahren zur Herstellung eines kalt- oder warmgewalzten Stahlbandes aus einem höchstfesten Mehrphasenstahl
DE102012006017A1 (de) 2012-03-20 2013-09-26 Salzgitter Flachstahl Gmbh Hochfester Mehrphasenstahl und Verfahren zur Herstellung eines Bandes aus diesem Stahl
DE102013004905A1 (de) 2012-03-23 2013-09-26 Salzgitter Flachstahl Gmbh Zunderarmer Vergütungsstahl und Verfahren zur Herstellung eines zunderarmen Bauteils aus diesem Stahl
WO2013144377A1 (fr) * 2012-03-30 2013-10-03 Voestalpine Stahl Gmbh Tôle d'acier laminée à froid à haute résistance et procédé de production d'une tôle d'acier de ce type
ES2746285T5 (es) * 2012-03-30 2022-12-19 Voestalpine Stahl Gmbh Lámina de acero de alta resistencia laminada en frío y procedimiento para producir dicha lámina de acero
ES2648415T5 (es) * 2012-03-30 2021-02-15 Voestalpine Stahl Gmbh Chapa de acero de alta resistencia laminada en frío y procedimiento de fabricación de dicha chapa de acero
MX366776B (es) * 2012-04-06 2019-07-23 Nippon Steel Corp Lamina de acero, laminada en caliente, recocida y galvanizada por inmersion en caliente, y proceso para producir la misma.
CA2869700C (fr) * 2012-04-12 2017-12-19 Jfe Steel Corporation Tole d'acier laminee a chaud pour tube d'acier carre destine a etre utilise comme element structural de construction et procede pour sa production
WO2013182622A1 (fr) * 2012-06-05 2013-12-12 Thyssenkrupp Steel Europe Ag Acier, produit en acier plat et procédé de fabrication d'un produit en acier plat
US10053757B2 (en) * 2012-08-03 2018-08-21 Tata Steel Ijmuiden Bv Process for producing hot-rolled steel strip
RU2507297C1 (ru) * 2012-10-05 2014-02-20 Леонид Михайлович Клейнер Стали со структурой пакетного мартенсита
EP2767601B1 (fr) * 2013-02-14 2018-10-10 ThyssenKrupp Steel Europe AG Produit plat en acier laminé à froid pour applications d'emboutissage profond et son procédé de fabrication
WO2014139625A1 (fr) * 2013-03-11 2014-09-18 Tata Steel Ijmuiden Bv Bande d'acier haute résistance, à phase complexe et galvanisée à chaud

Patent Citations (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
DE102012013113A1 (de) * 2012-06-22 2013-12-24 Salzgitter Flachstahl Gmbh Hochfester Mehrphasenstahl und Verfahren zur Herstellung eines Bandes aus diesem Stahl mit einer Mindestzugfestigkleit von 580MPa

Cited By (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP3825433B1 (fr) * 2018-08-22 2023-02-15 JFE Steel Corporation Tôle d'acier de haute résistance et méthode de production pour celle-ci
EP3950994B1 (fr) * 2019-03-28 2024-01-24 Nippon Steel Corporation Tôle d'acier à haute résistance

Also Published As

Publication number Publication date
MX2017006374A (es) 2018-02-16
US20190316222A1 (en) 2019-10-17
US10626478B2 (en) 2020-04-21
EP3221483A1 (fr) 2017-09-27
RU2017120860A (ru) 2018-12-19
RU2721767C2 (ru) 2020-05-22
DE102014017274A1 (de) 2016-05-19
CN107208232B (zh) 2019-02-26
WO2016078644A1 (fr) 2016-05-26
KR20170084210A (ko) 2017-07-19
RU2017120860A3 (fr) 2019-07-26
CN107208232A (zh) 2017-09-26

Similar Documents

Publication Publication Date Title
EP2864517B1 (fr) Acier multiphase à haute résistance et procédé pour la fabrication d'une bande faite de cet acier présentant une résistance à la traction minimale de 580 mpa
EP3027784B1 (fr) Acier multiphase à haute résistance, micro-allié et contenant du silicium, présentant une résistance minimale à la traction de 750 mpa et des propriétés améliorées et procédé de fabrication d'une bande à partir de cet acier
DE102012002079B4 (de) Verfahren zur Herstellung eines kalt- oder warmgewalzten Stahlbandes aus einem höchstfesten Mehrphasenstahl
EP3221483B1 (fr) Acier multiphases autotrempant en profondeur à haute résistance mécanique avec excellent aptitude au faconnage et procédé de fabrication d'une bande de cet acier
EP3221484B1 (fr) Procédé de production d'une bande en acier polyphasé, durcissant à l'air, ayant une haute résistance et ayant d'excellentes propriétés de mise en oeuvre
EP3221478B1 (fr) Bande à chaud ou à froid d'un acier multiphasé à haute résistance durcissant à l'air qui présente d'excellentes propriétés de traitement et procédé de fabrication d'une bande à chaud ou à froid à partir de cet acier multiphasé à haute résistance durcissant à l'air
EP2836614B1 (fr) Acier polyphasé à haute résistance et procédé de fabrication d'une bande à partir dudit acier
EP2855717B1 (fr) Tôle d'acier et méthode pour son obtention
EP2668302B1 (fr) Procédé de fabrication d'une bande d'acier à partir d'un acier multiphasé à haute résistance mécanique présentant d'excellentes propriétés de déformation
WO2017009192A1 (fr) Acier multiphase à haute résistance et procédé de fabrication d'une bande d'acier laminée à froid composée dudit acier
EP3692178B1 (fr) Procede de fabrication d'une bande d'acier a partir d'un acier multiphase a tres haute resistance
EP3964591A1 (fr) Produit en acier plat laminé à chaud et procédé de fabrication d'un produit en acier plat laminé à chaud
WO2022184811A1 (fr) Produit plat en acier, son procédé de production, et utilisation d'un tel produit plat en acier
EP3749469B1 (fr) Procédé de fabrication d'un élément structural par formage à chaud d'un produit primaire en acier manganésifère et élément structural formé à chaud
WO2024068957A1 (fr) Procédé de fabrication d'une bande d'acier à partir d'un acier multiphase à haute résistance et bande d'acier correspondante
EP4174207A1 (fr) Produit plat en acier ayant des propriétés de traitement améliorées

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: 20170607

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

Free format text: STATUS: EXAMINATION IS IN PROGRESS

17Q First examination report despatched

Effective date: 20180709

REG Reference to a national code

Ref country code: DE

Ref legal event code: R079

Ref document number: 502015012551

Country of ref document: DE

Free format text: PREVIOUS MAIN CLASS: C22C0038020000

Ipc: C21D0001840000

RIC1 Information provided on ipc code assigned before grant

Ipc: C21D 9/58 20060101ALI20190924BHEP

Ipc: C22C 38/22 20060101ALI20190924BHEP

Ipc: C22C 38/32 20060101ALI20190924BHEP

Ipc: C21D 9/56 20060101ALI20190924BHEP

Ipc: C21D 1/26 20060101ALI20190924BHEP

Ipc: C22C 38/06 20060101ALI20190924BHEP

Ipc: C21D 9/48 20060101ALI20190924BHEP

Ipc: C21D 8/04 20060101ALI20190924BHEP

Ipc: C23C 2/00 20060101ALI20190924BHEP

Ipc: C22C 38/28 20060101ALI20190924BHEP

Ipc: C21D 1/74 20060101ALI20190924BHEP

Ipc: C22C 38/02 20060101ALI20190924BHEP

Ipc: C21D 1/84 20060101AFI20190924BHEP

Ipc: C22C 38/04 20060101ALI20190924BHEP

Ipc: C22C 38/26 20060101ALI20190924BHEP

Ipc: C21D 1/28 20060101ALI20190924BHEP

Ipc: C22C 38/38 20060101ALI20190924BHEP

Ipc: C21D 1/76 20060101ALI20190924BHEP

Ipc: C21D 9/46 20060101ALI20190924BHEP

Ipc: C21D 8/02 20060101ALI20190924BHEP

GRAP Despatch of communication of intention to grant a patent

Free format text: ORIGINAL CODE: EPIDOSNIGR1

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

Free format text: STATUS: GRANT OF PATENT IS INTENDED

INTG Intention to grant announced

Effective date: 20191121

GRAS Grant fee paid

Free format text: ORIGINAL CODE: EPIDOSNIGR3

GRAA (expected) grant

Free format text: ORIGINAL CODE: 0009210

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

Free format text: STATUS: THE PATENT HAS BEEN GRANTED

AK Designated contracting states

Kind code of ref document: B1

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

REG Reference to a national code

Ref country code: GB

Ref legal event code: FG4D

Free format text: NOT ENGLISH

REG Reference to a national code

Ref country code: CH

Ref legal event code: EP

Ref country code: AT

Ref legal event code: REF

Ref document number: 1266803

Country of ref document: AT

Kind code of ref document: T

Effective date: 20200515

REG Reference to a national code

Ref country code: IE

Ref legal event code: FG4D

Free format text: LANGUAGE OF EP DOCUMENT: GERMAN

REG Reference to a national code

Ref country code: DE

Ref legal event code: R096

Ref document number: 502015012551

Country of ref document: DE

REG Reference to a national code

Ref country code: LT

Ref legal event code: MG4D

REG Reference to a national code

Ref country code: NL

Ref legal event code: MP

Effective date: 20200506

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: SE

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: LT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: NO

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200806

Ref country code: GR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200807

Ref country code: PT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200907

Ref country code: IS

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200906

Ref country code: FI

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: RS

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: HR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: LV

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: BG

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200806

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: NL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: AL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: RO

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: IT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: ES

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: DK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: SM

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: EE

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: CZ

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

REG Reference to a national code

Ref country code: DE

Ref legal event code: R097

Ref document number: 502015012551

Country of ref document: DE

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: SK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: PL

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

PLBE No opposition filed within time limit

Free format text: ORIGINAL CODE: 0009261

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

Free format text: STATUS: NO OPPOSITION FILED WITHIN TIME LIMIT

26N No opposition filed

Effective date: 20210209

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: SI

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: MC

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

REG Reference to a national code

Ref country code: CH

Ref legal event code: PL

GBPC Gb: european patent ceased through non-payment of renewal fee

Effective date: 20201106

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: LU

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201106

REG Reference to a national code

Ref country code: BE

Ref legal event code: MM

Effective date: 20201130

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: LI

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201130

Ref country code: CH

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201130

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: IE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201106

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: GB

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201106

REG Reference to a national code

Ref country code: AT

Ref legal event code: MM01

Ref document number: 1266803

Country of ref document: AT

Kind code of ref document: T

Effective date: 20201106

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: AT

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201106

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: TR

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: MT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

Ref country code: CY

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: MK

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20200506

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: BE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20201130

PGFP Annual fee paid to national office [announced via postgrant information from national office to epo]

Ref country code: FR

Payment date: 20231120

Year of fee payment: 9

Ref country code: DE

Payment date: 20231121

Year of fee payment: 9