EP3279353A1 - Tôle d'acier laminée à chaud et son procédé de production - Google Patents

Tôle d'acier laminée à chaud et son procédé de production Download PDF

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EP3279353A1
EP3279353A1 EP16771783.4A EP16771783A EP3279353A1 EP 3279353 A1 EP3279353 A1 EP 3279353A1 EP 16771783 A EP16771783 A EP 16771783A EP 3279353 A1 EP3279353 A1 EP 3279353A1
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steel sheet
phase
temperature
hot rolled
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German (de)
English (en)
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EP3279353A4 (fr
EP3279353B1 (fr
Inventor
Takaaki Tanaka
Taro Kizu
Shunsuke Toyoda
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JFE Steel Corp
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JFE Steel Corp
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    • 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
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
    • C—CHEMISTRY; METALLURGY
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
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    • 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
    • C21D6/02—Hardening by precipitation
    • 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
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • 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
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    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/10—Ferrous alloys, e.g. steel alloys containing cobalt
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/22—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
    • C22C38/28—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
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of 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
    • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/004—Dispersions; Precipitations
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—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

Definitions

  • the disclosure relates to a hot rolled steel sheet having high strength such as a tensile strength (TS) of 780 MPa or more, excellent stretch flangeability and blanking workability, and excellent manufacturing stability and suitable for structural-use steel material such as material for parts of transport machinery including vehicles and construction steel material.
  • TS tensile strength
  • the disclosure also relates to a method of manufacturing the hot rolled steel sheet.
  • An effective way of lightening automotive bodies while maintaining their strength is to strengthen steel sheets as material for automotive parts to thus reduce the thickness of steel sheets.
  • automotive suspension parts in which thick steel sheets tend to be used are expected to be lightened considerably by reducing the thickness of steel sheets through strengthening.
  • automotive suspension parts such as lower control arms are formed by burring, and so require steel sheets to have excellent stretch flangeability.
  • Much research and development have been conducted for hot rolled steel sheets having both strength and workability, and various techniques have been proposed. For example, it is known that high tensile strength and excellent stretch flangeability can both be achieved by making the metallic microstructure a substantially ferrite single-phase microstructure and precipitating fine carbides in the grains of the ferrite phase.
  • JP 2012-26034 A discloses a hot rolled steel sheet whose strength is improved while maintaining stretch flangeability, by making the steel sheet microstructure a ferrite single-phase microstructure having excellent workability with low dislocation density and dispersing and precipitating fine carbides in the ferrite to achieve strengthening by precipitation.
  • Burring is typically performed using a steel sheet blanked in a predetermined shape.
  • the part blanking clearance usually varies due to a temperature increase or wear of the tool caused by continuous pressing.
  • defects such as cracking or chipping may occur in the punched end surface. This has raised demand for a steel sheet that maintains excellent blanking workability regardless of the variations of the blanking conditions.
  • JP 2014-205888 A discloses a high strength hot rolled steel sheet whose mass-production blanking workability is improved by setting the volume fraction of bainite phase to more than 92%, setting the average spacing of bainite laths to 0.60 ⁇ m or less, and setting the number ratio of Fe-based carbides precipitated in the grains to all Fe-based carbides to 10% or more.
  • the steel sheet described in PTL 1 has both high strength and excellent stretch flangeability.
  • the steel sheet microstructure is a substantially ferrite single-phase microstructure, there is hardly any inclusion that serves as a void origin when blanking the steel sheet. Accordingly, in the steel sheet described in PTL 1, the punched end surface may become rough when conditions such as clearance and a blank holder vary.
  • the steel sheet described in PTL 2 has excellent blanking workability, by controlling the hot rolling conditions so that the steel sheet microstructure is mainly composed of predetermined bainite.
  • a bainite microstructure tends to vary in mechanical properties such as tensile strength due to variations in coiling temperature. It is often not easy to keep uniform steel sheet temperature throughout the length and width of the coil during cooling after hot rolling.
  • the steel sheet described in PTL 2 may thus vary greatly in mechanical properties, leading to lower manufacturing stability.
  • TS tensile strength
  • the steel sheet having the bainite single-phase microstructure has excellent stretch flangeability.
  • the steel sheet having the bainite single-phase microstructure also has excellent blanking workability, because many Fe-based carbides are present in the bainite microstructure and serve as a void origin during blanking.
  • the bainite microstructure varies greatly in mechanical properties such as strength depending on the transformation temperature, there is a possibility that the mechanical properties of the steel sheet vary greatly due to variations in hot rolling conditions such as coiling temperature.
  • Tempering a bainite or martensite microstructure typically enables a significant reduction of the variations of the mechanical properties caused by the variations of the hot rolling conditions, but also leads to a significant decrease in steel sheet strength. Besides, since Fe-based carbide morphology in tempered bainite or tempered martensite phase varies depending on the annealing conditions, the steel sheet may not be able to have excellent blanking workability depending on the annealing conditions.
  • the aforementioned steel sheet microstructure can be stably obtained particularly by adding 0.03% or more Ti and appropriately adjusting heat hysteresis in the annealing.
  • MC-type carbides are carbides, such as TiC, NbC, VC, and (Ti, Mo)C, with an atom ratio between an M element (for example, Ti, Nb, V, or Mo) and C of approximately 1:1.
  • M element for example, Ti, Nb, V, or Mo
  • the M element need not be of one type, and a complex carbide containing a plurality of metal elements is applicable.
  • a N-containing carbonitride or complex carbonitride is applicable, too.
  • a hot rolled steel sheet that has high strength such as a tensile strength (TS) of 780 MPa or more and excellent stretch flangeability and blanking workability and whose variations in mechanical properties caused by variations in manufacturing conditions are reduced, which is suitable for structural-use steel material such as material for parts of transport machinery including vehicles and construction steel material.
  • TS tensile strength
  • the C content improves the strength of the steel, and promotes the formation of bainite and martensite during hot rolling.
  • the C content therefore needs to be 0.03% or more. If the C content is more than 0.20%, equivalent carbon content is excessively high, which causes a decrease in weldability of the steel sheet.
  • the C content is therefore 0.03% or more and 0.20% or less.
  • the C content is preferably 0.04% or more.
  • the C content is preferably 0.18% or less.
  • the C content is more preferably more than 0.05%.
  • the C content is more preferably 0.15% or less.
  • Si is actively used in a high strength steel sheet as an effective element that improves the steel sheet strength without decreasing ductility (elongation). If the Si content is more than 0.4%, however, Si forms oxides on the steel sheet surface during heat treatment, and degrades coating adhesion property. The Si content is therefore 0.4% or less. The Si content is preferably 0.3% or less. The Si content is more preferably 0.2% or less. The Si content may be reduced to an impurity level, and may be 0%.
  • Mn 0.5% or more and 2.0% or less
  • Mn is an element that dissolves and contributes to higher strength of the steel. Mn also promotes the formation of bainite and martensite during hot rolling, by improving quench hardenability. To achieve such effects, the Mn content needs to be 0.5% or more. If the Mn content is more than 2.0%, austenite becomes excessively stable, causing the microstructure of the steel sheet to excessively contain martensite and retained austenite. This decreases stretch flangeability. The Mn content is therefore 0.5% or more and 2.0% or less. The Mn content is preferably 0.8% or more. The Mn content is preferably 1.8% or less. The Mn content is more preferably 1.0% or more. The Mn content is more preferably 1.7% or less.
  • P is a harmful element that segregates to grain boundaries to decrease elongation, induce cracking during working, and degrade anti-crash property.
  • the P content is therefore 0.03% or less. Excessive dephosphorization, however, leads to longer refining time and higher cost, and so the P content is preferably 0.002% or more.
  • S exists as MnS or TiS in the steel, and facilitates the formation of voids when blanking the hot rolled steel sheet. S also serves as a void origin during working, and causes a decrease in stretch flangeability.
  • the S content is therefore desirably as low as possible, and is 0.03% or less.
  • the S content is preferably 0.01% or less. Excessive desulfurization, however, leads to longer refining time and higher cost, and so the S content is preferably 0.0002% or more.
  • Al is an element that acts as a deoxidizing material.
  • the Al content is desirably 0.01% or more. If the Al content is more than 0.1%, Al remains in the steel sheet as Al oxide. Such Al oxide tends to coagulate and be coarsened, causing a decrease in stretch flangeability.
  • the Al content is therefore 0.1% or less.
  • N exists as coarse TiN in the steel, and facilitates the formation of coarse voids when blanking the hot rolled steel sheet. N also serves as an origin of coarse voids during working, and causes a decrease in stretch flangeability.
  • the N content is therefore desirably as low as possible, and is 0.01% or less.
  • the N content is preferably 0.006% or less. Excessive denitrification, however, leads to longer refining time and higher cost, and so the N content is preferably 0.0005% or more.
  • Ti is a necessary element to form MC-type carbides to thus inhibit lath coarsening in the annealing and strengthen the steel sheet.
  • MC-type carbides also enhance the steel sheet strength by strengthening by precipitation. If the Ti content is less than 0.03%, such effects are insufficient, and lath coarsening and lower precipitation amount cause a decrease in steel sheet strength, making it difficult to achieve desired steel sheet strength (tensile strength of 780 MPa or more). If the Ti content is more than 0.15%, central segregation is noticeable, causing a decrease in blanking workability.
  • the Ti content is therefore 0.03% or more and 0.15% or less.
  • the Ti content is preferably 0.04% or more.
  • the Ti content is preferably 0.14% or less.
  • the Ti content is further preferably 0.05% or more.
  • the Ti content is further preferably 0.13% or less.
  • the hot rolled steel sheet may optionally contain one or more of V: 0.01% or more and 0.3% or less, Nb: 0.01% or more and 0.1% or less, and Mo: 0.01% or more and 0.3% or less, for higher strength.
  • V 0.01% or more and 0.3% or less
  • V forms MC-type carbides and contributes to higher strength of the steel sheet by lath coarsening inhibition in the annealing and strengthening by precipitation, as with Ti. To achieve such effects, the V content needs to be 0.01% or more. If the V content is more than 0.3%, central segregation is noticeable, causing a decrease in blanking workability. Accordingly, the V content is preferably 0.01% or more. The V content is preferably 0.3% or less. The V content is more preferably 0.01% or more. The V content is more preferably 0.2% or less. The V content is further preferably 0.01% or more. The V content is further preferably 0.15% or less. V may form MC-type carbides by itself, or form complex carbides with Ti, Nb, and Mo. Such carbide composition does not affect the advantageous effects of the disclosure at all.
  • Nb 0.01% or more and 0.1% or less
  • Nb forms MC-type carbides and contributes to higher strength of the steel sheet by lath coarsening inhibition in the annealing and strengthening by precipitation, as with Ti.
  • the Nb content needs to be 0.01% or more. If Nb is excessively added to be more than 0.1% in content, Nb does not dissolve in the heating furnace during hot rolling. The effects thus saturate, and the alloy cost increases. Accordingly, the Nb content is preferably 0.01% or more.
  • the Nb content is preferably 0.1 % or less.
  • the Nb content is more preferably 0.01% or more.
  • the Nb content is more preferably 0.08% or less.
  • the Nb content is further preferably 0.01% or more.
  • the Nb content is further preferably 0.06% or less.
  • Nb may form MC-type carbides by itself, or form complex carbides with Ti, V, and Mo. Such carbide composition does not affect the advantageous effects of the disclosure at all.
  • Mo when added in combination with Ti, forms MC-type complex carbides and contributes to higher strength of the steel sheet by lath coarsening inhibition in the annealing and strengthening by precipitation, as with Ti.
  • the Mo content needs to be 0.01% or more. If the Mo content is more than 0.3%, central segregation is noticeable, causing a decrease in blanking workability. Accordingly, the Mo content is preferably 0.01% or more. The Mo content is preferably 0.3% or less. Mo may form complex carbides with Nb and V. Such carbide composition does not affect the advantageous effects of the disclosure at all.
  • the hot rolled steel sheet may optionally contain B: 0.0002% or more and 0.010% or less, for improved quench hardenability during hot rolling.
  • the B is an element that segregates to austenite grain boundaries and inhibits the formation and growth of ferrite to improve quench hardenability and promote the formation of bainite and martensite.
  • the B content is preferably 0.0002% or more. If the B content is more than 0.010%, hard iron boride forms and causes a decrease in stretch flangeability. Accordingly, in the case of adding B, the B content is preferably 0.0002% or more and 0.010% or less.
  • the B content is more preferably 0.0002% or more.
  • the B content is more preferably 0.0050% or less.
  • the B content is further preferably 0.0004% or more.
  • the B content is further preferably 0.0030% or less.
  • the hot rolled steel sheet may contain one or more of REM, Zr, As, Cu, Ni, Sn, Pb, Ta, W, Cr, Sb, Mg, Ca, Co, Se, Zn, and Cs so that their total content is 1.0% or less.
  • the components other than those described above are Fe and incidental impurities.
  • Total area ratio of tempered bainite phase and tempered martensite phase 70% or more
  • the hot rolled steel sheet has a microstructure mainly composed of tempered bainite and tempered martensite having both high strength and excellent blanking workability. If the total area ratio of tempered bainite phase and tempered martensite phase is less than 70%, the hot rolled steel sheet cannot have desired high strength and blanking workability.
  • the ratio of each of tempered bainite phase and tempered martensite is not individually defined because tempered bainite and tempered martensite after annealing are microstructures not distinguishable from each other. This is a major factor that can reduce variations in mechanical properties after annealing in the case where the manufacturing conditions during hot rolling vary.
  • the total area ratio of tempered bainite phase and tempered martensite phase is therefore 70% or more.
  • the total area ratio of tempered bainite phase and tempered martensite phase is preferably 75% or more.
  • the total area ratio of tempered bainite phase and tempered martensite phase is more preferably 80% or more.
  • the total area ratio of tempered bainite phase and tempered martensite phase may be 100%.
  • Total area ratio of coarse pearlite phase, martensite phase, and retained austenite phase 10% or less
  • the microstructure of the hot rolled steel sheet is mainly composed of tempered bainite and tempered martensite, with the balance other than tempered bainite and tempered martensite being, for example, Fe-based carbides, coarse pearlite, fine pearlite, degenerate pearlite, bainite, martensite, and retained austenite.
  • tempered bainite and tempered martensite being, for example, Fe-based carbides, coarse pearlite, fine pearlite, degenerate pearlite, bainite, martensite, and retained austenite.
  • the total area ratio of coarse pearlite phase, martensite phase, and retained austenite phase is therefore 10% or less.
  • the total area ratio of coarse pearlite phase, martensite phase, and retained austenite phase is preferably 8% or less.
  • the total area ratio of coarse pearlite phase, martensite phase, and retained austenite phase is more preferably 5% or less.
  • coarse pearlite has a lamellar spacing of 0.2 ⁇ m or more
  • fine pearlite has a lamellar spacing of less than 0.2 ⁇ m
  • degenerate pearlite is a phase in which pearlite lamellar is not clearly observable.
  • the lamellar spacing can be measured by microstructure observation using a scanning electron microscope.
  • the balance other than tempered bainite phase, tempered martensite phase, coarse pearlite phase, martensite phase, and retained austenite phase is, for example, ferrite phase, degenerate pearlite phase, and fine pearlite phase.
  • a total area ratio of such balance of 30% or less is allowable.
  • Average width of laths which tempered bainite phase and tempered martensite phase have as substructure 1.0 ⁇ m or less
  • FIG. 1 is a schematic diagram illustrating an example of a microstructure in which tempered bainite phase and tempered martensite phase have laths as their substructure and Fe-based carbides precipitate and MC-type carbides disperse and precipitate inside and at the boundaries of the laths. If the laths disappears as a result of recovery or the average width of the laths is more than 1.0 ⁇ m, predetermined high strength cannot be achieved.
  • the average width of laths which tempered bainite phase and tempered martensite phase have as their substructure is therefore 1.0 ⁇ m or less.
  • the average width of laths is preferably 0.8 ⁇ m or less.
  • the average width of laths is more preferably 0.6 ⁇ m or less. No lower limit is placed on the average width of laths, yet the lower limit is typically about 0.1 ⁇ m.
  • Fe-based carbides precipitated inside and at the boundaries of laths as illustrated in FIG. 1 serve as a void origin during blanking, thus contributing to improved blanking workability. This effect is particularly high with Fe-based carbides having an aspect ratio of 5 or less.
  • the proportion of Fe-based carbides with an aspect ratio of 5 or less in Fe-based carbides precipitated inside and at the boundaries of laths is therefore 80% or more.
  • the proportion is preferably 85% or more. No upper limit is placed on the proportion, yet the upper limit may be 100%.
  • Fe-based carbides are ⁇ carbide (cementite), ⁇ carbide, and the like.
  • An alloying element may be dissolved in the carbides.
  • the aspect ratio is the ratio of the major axis length and minor axis length of Fe-based carbides precipitated inside and at the boundaries of laths.
  • Average particle size of MC-type carbides dispersed and precipitated inside and at the boundaries of laths 20 nm or less
  • MC-type carbides finely dispersed and precipitated inside and at the boundaries of laths as illustrated in FIG. 1 inhibit lath coarsening by a pinning effect when annealing the steel sheet and also inhibit lath disappearance resulting from recovery, thus contributing to higher strength. If the average particle size of MC-type carbides is more than 20 nm, the number of particles of MC-type carbides contributing to pinning is insufficient and so the pinning effect is insufficient, causing a decrease in steel sheet strength. If the average particle size of MC-type carbides is 20 nm or less, a sufficient number of particles of MC-type carbides exhibit the pinning effect, to prevent a decrease in steel sheet strength.
  • the average particle size of MC-type carbides dispersed and precipitated inside and at the boundaries of laths of tempered bainite phase and tempered martensite phase is therefore 20 nm or less.
  • the average particle size is preferably 15 nm or less. No lower limit is placed on the average particle size, yet the lower limit is typically about 1 nm.
  • the proportion of MC-type carbides with a particle size of more than 50 nm is preferably 10% or less.
  • the hot rolled steel sheet has an average dislocation density in the following range.
  • Average dislocation density 1.0 ⁇ 10 14 m -2 or more and 5.0 ⁇ 10 15 m -2 or less
  • the variations of the hot rolled steel sheet caused by the variations of the hot rolling conditions are reduced by tempering the steel sheet having bainite and martensite microstructure. If the average dislocation density of the steel sheet after annealing is more than 5.0 ⁇ 10 15 m -2 , the tempering of the steel sheet is insufficient and the influence of the variations of the hot rolling conditions cannot be reduced sufficiently. In the case where tempering is sufficient, the average dislocation density is typically 1.0 ⁇ 10 14 m -2 or more. The average dislocation density is therefore 1.0 ⁇ 10 14 m -2 or more and 5.0 ⁇ 10 15 m -2 or less. The average dislocation density is preferably 1.0 ⁇ 10 14 m -2 or more. The average dislocation density is preferably 2.0 ⁇ 10 15 m -2 or less.
  • the method of manufacturing the hot rolled steel sheet includes: hot rolling a steel raw material having the chemical composition described above, whereby the steel raw material is heated to an austenite single phase region, subjected to rough rolling and finish rolling to obtain a steel sheet, and the steel sheet is cooled and coiled after the finish rolling; pickling the steel sheet after the hot rolling; and then continuous annealing the steel sheet, wherein in the hot rolling, a finisher delivery temperature is 850 °C or more and 1000 °C or less, an average cooling rate to 500 °C after the finish rolling is 30 °C/s or more, and a coiling temperature is 500 °C or less, and in the continuous annealing, a maximum heating temperature of the steel sheet is 700 °C or more and (A 3 point + A 1 point)/2 or less, a time during which a temperature of the steel sheet is 600 °C or more and 700 °C or less in heating the steel sheet to the maximum heating temperature is 20 s or more and 1000 s or less, a time during
  • the method of obtaining the steel raw material by steelmaking is not limited, and any known steelmaking process such as a converter steelmaking process or an electric furnace steelmaking process may be used.
  • continuous casting is preferably performed to yield a slab (steel raw material) in terms of productivity and the like.
  • the slab may be yielded by a known casting method such as ingot casting and blooming, thin slab continuous casting, or the like.
  • the obtained steel raw material is subjected to hot rolling, in which the steel raw material is subjected to rough rolling and finish rolling.
  • the steel raw material is heated in the austenite single phase region. If the steel raw material before the rough rolling is not heated in the austenite single phase region, the remelting of Ti carbide and the like present in the steel raw material does not progress, and fine MC-type carbides do not precipitate during the annealing after the hot rolling. Accordingly, the steel raw material is heated to the austenite single phase region, preferably to 1150 °C or more, before the rough rolling.
  • the heating temperature is typically 1350 °C or less.
  • the steel raw material may be subjected to hot direct rolling without being heated or after being heated for a short time.
  • Finisher delivery temperature 850 °C or more and 1000 °C or less
  • the finisher delivery temperature needs to be 850 °C or more.
  • the finisher delivery temperature is preferably 880 °C or more. If the finisher delivery temperature is more than 1000 °C, the surface characteristics of the steel sheet degrade. The finisher delivery temperature is therefore 1000 °C or less.
  • the finisher delivery temperature is preferably 970 °C or less.
  • Each of the temperatures such as the finisher delivery temperature and the coiling temperature mentioned here is the temperature of the steel sheet surface.
  • Cooling rate to 500 °C after finish rolling 30 °C/s or more
  • the cooling rate to 500 °C after the finish rolling needs to be 30 °C/s or more.
  • the cooling rate is preferably 50 °C/s or more. No upper limit is placed on the cooling rate, yet the upper limit is typically about 300 °C/s.
  • Coiling temperature 500 °C or less
  • the coiling temperature is important in controlling the steel sheet microstructure after the hot rolling. If the coiling temperature is more than 500 °C, the lath width of bainite increases. This makes it impossible to obtain a predetermined lath width of tempered bainite after the annealing. No lower limit is placed on the coiling temperature, yet an excessively low coiling temperature merely leads to higher cooling cost, and so the coiling temperature is preferably 0 °C or more. The coiling temperature is more preferably 200 °C or more.
  • the hot rolled steel sheet After the hot rolling, the hot rolled steel sheet is subjected to pickling and then to continuous annealing.
  • the reasons for limiting the manufacturing conditions in the continuous annealing are given below.
  • Appropriately adjusting the maximum heating temperature of the steel sheet in the continuous annealing is important in sufficiently reducing the influence of the variations of the manufacturing conditions in the hot rolling caused by the annealing and achieving desired high strength. If the maximum heating temperature of the steel sheet is less than 700 °C, the dislocation density in bainite and martensite is difficult to be controlled within an appropriate range, and so the influence of the variations of the manufacturing conditions in the hot rolling cannot be reduced sufficiently. Besides, if the heating temperature of the steel sheet is less than 700 °C, the aspect ratio of Fe-based carbides inside and between laths tends to be high, which makes it difficult to set the proportion of Fe-based carbides with an aspect ratio of 5 or less to be in a desired range.
  • the maximum heating temperature of the steel sheet in the continuous annealing is therefore 700 °C or more and (A 3 point + A 1 point)/2 or less.
  • the maximum heating temperature is preferably 700 °C or more.
  • the maximum heating temperature is preferably ⁇ (A 3 point + A 1 point)/2 ⁇ - 10 °C or less.
  • the A 1 point and the A 3 point can be calculated according to the following expressions.
  • a 1 point 751 ⁇ 26.6 ⁇ % C + 17.6 ⁇ % Si ⁇ 11.6 ⁇ % Mn + 22.5 ⁇ % Mo + 233 ⁇ % Nb ⁇ 39.7 ⁇ % V ⁇ 57 ⁇ % Ti ⁇ 895 ⁇ % B ⁇ 169 ⁇ % Al
  • a 3 point 937 ⁇ 476.5 ⁇ % C + 56 ⁇ % Si ⁇ 19.7 ⁇ % Mn + 38.1 ⁇ % Mo + 124.8 ⁇ % V + 136.3 ⁇ % Ti ⁇ 19 ⁇ % Nb + 3315 ⁇ % B
  • [%X] denotes the content of an X element in steel (mass%).
  • MC-type carbides In heating the steel sheet to the maximum heating temperature, it is important to appropriately control heat hysteresis in imparting desired high strength and excellent blanking workability to the steel sheet.
  • the pinning effect of MC-type carbides is used to inhibit lath coarsening, as mentioned above. To achieve the pinning effect, MC-type carbides need to be sufficiently dispersed in bainite and martensite before lath coarsening starts. According to our study, the precipitation of MC-type carbides begins to occur noticeably at 600 °C or more. Meanwhile, lath coarsening and disappearance are noticeable at more than 700 °C.
  • lath coarsening and disappearance can be inhibited by holding the steel sheet temperature in the temperature range of 600 °C or more and 700 °C or less for a predetermined time so that MC-type carbides precipitate sufficiently.
  • the holding time in this temperature range needs to be 20 s or more. If the holding time in the temperature range is insufficient, lath coarsening starts before MC-type carbides precipitate sufficiently, so that the pinning effect is insufficient and the laths coarsen.
  • the holding time is preferably 35 s or more.
  • the holding time is more preferably 50 s or more.
  • the holding time of the steel sheet temperature in the temperature range of 600 °C or more and 700 °C or less is more than 1000 s, Fe-based carbides precipitated inside and between laths dissolve again and move to prior austenite grain boundaries, packet grain boundaries, block grain boundaries, and the like. Thus, Fe-based carbides inside and between laths that effectively contribute to improved blanking workability no longer exist. Accordingly, to obtain a steel sheet having excellent blanking workability, the holding time of the steel sheet temperature in the temperature range of 600 °C or more and 700 °C or less needs to be 1000 s or less.
  • the holding time is preferably 800 s or less.
  • the holding time is more preferably 500 s or less.
  • the steel sheet temperature mentioned here is the temperature of the steel sheet surface.
  • the holding time of the steel sheet temperature in the temperature range of more than 700 °C is 200 s or less, in terms of preventing lath coarsening.
  • the holding time is preferably 180 s or less.
  • the holding time is more preferably 150 s or less. If the time during which the steel sheet temperature is more than 700 °C is less than 10 s, the ductility of the steel sheet decreases to some extent, and so the holding time is preferably 10 s or more.
  • Average cooling rate to 530 °C when cooling the steel sheet from the maximum heating temperature 8 °C/s or more and 25 °C/s or less
  • non-transformed austenite transforms to martensite or remains in the steel sheet microstructure as retained austenite, so that stretch flangeability decreases.
  • the holding time in the temperature range of 470 °C or more and 530 °C or less after the controlled cooling stops is 10 s or more.
  • the holding time is preferably 20 s or more.
  • the holding time is more preferably 30 s or more. No upper limit is placed on the holding time, yet the holding time is typically 300 s or less.
  • Holding the steel sheet in the temperature range of 470 °C or more and 530 °C or less completes the control of the steel sheet microstructure.
  • the subsequent cooling conditions are not limited, and the steel sheet may be cooled to room temperature by any cooling method.
  • desired steel sheet microstructure can still be obtained as long as the total holding time in the temperature range of 600 °C or more and 700 °C or less is 1000 s or less.
  • the steel sheet may be immersed in a zinc pot to yield a hot-dip galvanized steel sheet.
  • the steel sheet may then be further heated to yield a galvannealed steel sheet.
  • the hot dip coating is not limited to zinc, and may be a coating of aluminum, an aluminum alloy, or the like.
  • the steel sheet may be subjected to temper rolling either continuously in the annealing line or using another line according to a conventional method.
  • the hot rolled steel sheet manufactured as described above may be electrogalvanized or hot-dip galvanized.
  • the hot rolled steel sheet according to the disclosure is suitable not only as a steel sheet for automotive suspension parts but also for press forming typically performed at ordinary temperature, and has excellent heat resistance.
  • the hot rolled steel sheet manufactured as described above is also suitable as a blank sheet for a warm forming process of heating a steel sheet to 400 °C to 700 °C before pressing and then immediately press forming the steel sheet.
  • Molten steels having the compositions listed in Table 1 were each obtained by steelmaking and subjected to continuous casting by a typically known technique, to yield a slab (steel raw material) with a thickness of 300 mm.
  • the slab was heated to the temperature in Table 2, rough rolled, and finish rolled at the finisher delivery temperature in Table 2.
  • the steel sheet was cooled at the average cooling rate in Table 2, and coiled at the coiling temperature in Table 2, to obtain a hot rolled steel sheet with a sheet thickness of 3.2 mm.
  • the hot rolled steel sheet was then pickled by a typically known technique, and annealed in a continuous annealing line under the conditions in Table 2.
  • test piece was collected from each obtained hot rolled steel sheet, and subjected to microstructure observation, average dislocation density measurement, a tensile test, a hole expansion test, a blanking test, and manufacturing stability evaluation.
  • the evaluation results are listed in Table 3.
  • the test methods are as follows.
  • a test piece was collected from each obtained hot rolled steel sheet, and polished in a cross-section (L cross-section) parallel to the rolling direction of the test piece and etched by nital.
  • a micrograph taken with a scanning electron microscope 1000, 3000, 5000 magnifications was used to determine the total area ratio of tempered bainite phase and tempered martensite phase, the area ratio of coarse pearlite phase, the total area ratio of martensite phase and retained austenite phase (MA), and the area ratio of phase other than these, through the use of an image analyzer. It is difficult to distinguish martensite phase and retained austenite phase from each other with a scanning electron micrograph.
  • the total area ratio of coarse pearlite phase, martensite phase, and retained austenite phase is important, and accordingly the total area ratio of martensite phase and retained austenite phase (MA) was determined without distinguishing martensite phase and retained austenite phase from each other.
  • a thin film made from each hot rolled steel sheet was observed using a transmission electron microscope (TEM), to measure the lath width in tempered bainite and tempered martensite and determine the proportion of Fe-based carbides with an aspect ratio of 5 or less in Fe-based carbides precipitated inside and at the boundaries of laths and the average particle size of MC-type carbides precipitated inside and at the boundaries of laths.
  • TEM transmission electron microscope
  • the lath width in tempered bainite and tempered martensite was measured as follows. In a transmission electron micrograph of 120 mm ⁇ 80 mm in size taken for 10 observation fields at 30000 magnifications, five straight lines orthogonal to the major axes of three or more consecutively aligned laths were drawn at intervals of 10 mm, the length of each line segment where the corresponding straight line intersects with the lath boundaries was measured, and the average length of the line segments was set as the average lath width.
  • the proportion of Fe-based carbides with an aspect ratio of 5 or less in Fe-based carbides precipitated inside and at the boundaries of laths was determined as follows. In a micrograph taken at 165000 magnifications, the major axis length and the minor axis length were measured for at least 100 particles of Fe-based carbides precipitated inside and at the boundaries of laths for 5 observation fields in total, to calculate the aspect ratio. The proportion of Fe-based carbides with an aspect ratio of 5 or less was thus determined.
  • the average particle size of MC-type carbides was determined as follows. In a micrograph taken at 300000 magnifications, the diameter was measured for at least 100 particles of MC-type carbides such as TiC for 5 observation fields in total, and an arithmetic average (average particle size d def ) was calculated. The lower limit of the measured particle size was 2 nm.
  • a test piece was collected from each obtained hot rolled steel sheet, and the dislocation density of a 1/4 portion in sheet thickness was measured. Assuming that the dislocation density of a 1/4 portion in sheet thickness represents the average dislocation density of the steel sheet, the measurement was set as average dislocation density.
  • the collected test piece was subjected to mechanical grinding and also polishing with oxalic acid for 0.1 mm, to adjust the sample so that the 1/4 portion in sheet thickness was exposed to the surface. Polishing with oxalic acid was intended to remove the layer worked by grinding.
  • the strain of the steel sheet was measured by an X-ray diffractometer.
  • an X-ray diffractometer With an X-ray diffractometer, the diffraction intensity of (110) plane, (211) plane, and (220) plane of ⁇ -iron in the 1/4 portion in sheet thickness was measured using CoK ⁇ rays.
  • the half-value breadth of the peak value of the reflection intensity of each crystal plane was calculated from the obtained measurement chart, and the local strain ⁇ ' applied to the steel sheet was determined according to the following Expressions (1) and (2).
  • ⁇ the half-value breadth of the peak value (the value corrected according to Expression (2) was used)
  • ⁇ the diffraction angle
  • ⁇ the wavelength of CoK ⁇ rays (0.1790 nm)
  • D the crystallite size (dislocation cell, crystal grain size)
  • ⁇ ' the local strain.
  • ⁇ 2 ⁇ m 2 ⁇ ⁇ 0 2
  • ⁇ m the half-value breadth of the peak of the sample subjected to dislocation density measurement
  • ⁇ 0 the half-value breadth of the peak of a strain-free sample.
  • JIS Z 2001 A JIS No. 5 tensile test piece (JIS Z 2001) was collected from each obtained hot rolled steel sheet so that the direction (C direction) orthogonal to the rolling direction was the tensile direction, and subjected to a tensile test in conformity with JIS Z 2241 to measure yield strength (YS), tensile strength (TS), and elongation (El).
  • test piece size: 100 mm ⁇ 100 mm
  • test piece blanked with a hole of 10 mm ⁇ in initial diameter do (clearance: 12.5% of the test piece sheet thickness).
  • a hole expansion test was conducted using the test piece.
  • a conical punch with a vertex angle of 60° was inserted into the hole of 10 mm ⁇ in initial diameter d 0 from the punch side at the time of blanking, to expand the hole.
  • the diameter d (mm) of the hole when a crack ran through the steel sheet (test piece) was measured, and the hole expansion ratio ⁇ (%) was calculated according to the following expression.
  • Hole expansion ratio ⁇ % d ⁇ d 0 / d 0 ⁇ 100.
  • the stretch flangeability was evaluated as favorable in the case where tensile strength (TS) ⁇ ⁇ hole expansion ratio ( ⁇ ) ⁇ 0.5 was 6200 ⁇ MPa% 0,5 or more.
  • test piece (size: 30 mm ⁇ 30 mm) was collected from each obtained hot rolled steel sheet, and blanked with a hole of 10 mm ⁇ in diameter do (clearance: 20%, 30% of the test piece sheet thickness). After the blanking, the fracture state of the punched end surface was observed by a microscope (50 magnifications) on the whole circumference of the punch hole, to observe whether or not any crack, chip, or brittle fracture occurred. The blanking workability was evaluated as "pass” if there was no crack, chip, or brittle fracture, and "fail” otherwise. Table 3 Steel sheet No. Steel No.
  • JIS No. 5 tensile test pieces JIS Z 2001 were optionally collected from the whole length and whole width of the hot rolled steel sheets of Examples so that the orthogonal direction (C direction) was the tensile direction.
  • C direction the orthogonal direction
  • TS tensile strength
  • ⁇ the standard deviation of the tensile strength
  • the mechanical properties of the steel sheet such as tensile strength (TS) had little variations, exhibiting excellent manufacturing stability.

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JP5728836B2 (ja) * 2009-06-24 2015-06-03 Ｊｆｅスチール株式会社 耐硫化物応力割れ性に優れた油井用高強度継目無鋼管の製造方法
JP5765080B2 (ja) * 2010-06-25 2015-08-19 Ｊｆｅスチール株式会社 伸びフランジ性に優れた高強度熱延鋼板およびその製造方法
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JP5126326B2 (ja) 2010-09-17 2013-01-23 Ｊｆｅスチール株式会社 耐疲労特性に優れた高強度熱延鋼板およびその製造方法
JP5724267B2 (ja) * 2010-09-17 2015-05-27 Ｊｆｅスチール株式会社 打抜き加工性に優れた高強度熱延鋼板およびその製造方法
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JP5454745B2 (ja) * 2011-10-04 2014-03-26 Ｊｆｅスチール株式会社 高強度鋼板およびその製造方法
JP5541263B2 (ja) * 2011-11-04 2014-07-09 Ｊｆｅスチール株式会社 加工性に優れた高強度熱延鋼板およびその製造方法
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JP5641087B2 (ja) 2013-04-15 2014-12-17 Ｊｆｅスチール株式会社 量産打抜き性に優れた高強度熱延鋼板およびその製造方法
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US11965223B2 (en) 2018-07-31 2024-04-23 Jfe Steel Corporation Thin steel sheet and method for manufacturing the same
EP4083241A4 (fr) * 2019-12-23 2023-08-16 Nippon Steel Corporation Tôle d'acier laminée à chaud

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MX2017012493A (es) 2018-01-18
CN107429362B (zh) 2020-06-23
WO2016157896A1 (fr) 2016-10-06
US20180119240A1 (en) 2018-05-03
EP3279353A4 (fr) 2018-02-07
JP6075517B1 (ja) 2017-02-08
EP3279353B1 (fr) 2019-03-27
KR101989262B1 (ko) 2019-06-13
JPWO2016157896A1 (ja) 2017-04-27
CN107429362A (zh) 2017-12-01
KR20170128555A (ko) 2017-11-22

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