WO2016190370A1 - Tôle d'acier et son procédé de production - Google Patents

Tôle d'acier et son procédé de production Download PDF

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WO2016190370A1
WO2016190370A1 PCT/JP2016/065509 JP2016065509W WO2016190370A1 WO 2016190370 A1 WO2016190370 A1 WO 2016190370A1 JP 2016065509 W JP2016065509 W JP 2016065509W WO 2016190370 A1 WO2016190370 A1 WO 2016190370A1
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carbides
steel sheet
annealing
temperature
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PCT/JP2016/065509
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English (en)
Japanese (ja)
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健悟 竹田
匹田　和夫
高田　健
元仙 橋本
友清　寿雅
保嗣 塚野
荒牧　高志
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新日鐵住金株式会社
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Priority to EP16800074.3A priority Critical patent/EP3305929A4/fr
Priority to US15/576,177 priority patent/US20180230582A1/en
Priority to BR112017024692-9A priority patent/BR112017024692A2/pt
Priority to MX2017015016A priority patent/MX2017015016A/es
Priority to JP2016559466A priority patent/JP6119923B1/ja
Priority to CN201680030099.7A priority patent/CN107614727B/zh
Priority to KR1020177033291A priority patent/KR102029565B1/ko
Publication of WO2016190370A1 publication Critical patent/WO2016190370A1/fr

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    • 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
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    • 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
    • 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
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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
    • 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
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    • 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
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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
    • 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
    • C21D8/0463—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 following hot rolling
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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
    • 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
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    • 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
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    • 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
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    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
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    • 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
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
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    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
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    • 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
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    • 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
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
    • C23G—CLEANING OR DE-GREASING OF METALLIC MATERIAL BY CHEMICAL METHODS OTHER THAN ELECTROLYSIS
    • C23G1/00—Cleaning or pickling metallic material with solutions or molten salts
    • C23G1/02—Cleaning or pickling metallic material with solutions or molten salts with acid solutions
    • C23G1/08—Iron or steel
    • 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

Definitions

  • the present invention relates to a steel plate and a manufacturing method thereof.
  • Cold forging is applied, and it is used as a material for driving system parts such as automobile gears and clutches.
  • cold forging increases the amount of strain accumulated in the material, causing cracks in the material and buckling at the time of molding, causing a problem of deterioration of component characteristics.
  • Patent Documents 1 to 5 So far, many proposals have been made on techniques for improving the cold forgeability of steel sheets and the impact resistance after carburizing (see, for example, Patent Documents 1 to 5).
  • Patent Document 1 as steel for mechanical structure whose toughness is improved by suppressing coarsening of crystal grains in carburizing heat treatment, C: 0.10 to 0.30%, Si: 0.05 ⁇ 2.0%, Mn: 0.10 ⁇ 0.50%, P: 0.030% or less, S: 0.030% or less, Cr: 1.80 ⁇ 3.00%, Al: 0.005 ⁇ 0.05%, Nb: 0.02 to 0.10%, N: 0.0300% or less, the balance is Fe and inevitable impurities, the structure before cold working is a ferrite pearlite structure, A steel for machine structural use having an average ferrite grain size of 15 ⁇ m or more is disclosed.
  • Patent Document 2 as steel having excellent cold workability and carburizing hardenability, C: 0.15 to 0.40%, Si: 1.00% or less, Mn: 0.40% or less, sol. Al: 0.02% or less, N: 0.006% or less, B: 0.005 to 0.050%, with the balance being Fe and inevitable impurities, mainly composed of ferrite phase and graphite phase A steel having a structure that achieves this is disclosed.
  • Patent Document 3 discloses a steel material for a carburized bevel gear having excellent impact strength, a high toughness carburized bevel gear, and a manufacturing method thereof.
  • Patent Document 4 cold forging is performed after spheroidizing annealing, and the parts produced in the carburizing and quenching and tempering process have excellent workability while suppressing the coarsening of crystal grains even in subsequent carburizing.
  • steel for carburized parts having excellent impact resistance characteristics and impact fatigue resistance characteristics is disclosed.
  • Patent Document 5 as cold tool steel for plasma carburizing, C: 0.40 to 0.80%, Si: 0.05 to 1.50%, Mn: 0.05 to 1.50%, and V: 1.8 to 6.0%, Ni: 0.10 to 2.50%, Cr: 0.1 to 2.0%, and Mo: 3.0% or less
  • the steel which contains 2 or more types and remainder consists of Fe and an unavoidable impurity is disclosed.
  • JP 2013-040376 A Japanese Patent Laid-Open No. 06-116679 JP 09-201644 A JP 2006-213951 A Japanese Patent Laid-Open No. 10-158780
  • the structure of the steel for machine structural use in Patent Document 1 is a structure of ferrite + pearlite, which has a larger hardness than the structure of ferrite + cementite, and therefore suppresses wear of the mold in cold forging. It cannot be said that it is necessarily a steel for machine structural use with excellent cold forgeability.
  • Patent Document 3 is not a manufacturing method that drastically lowers the cost because hot forging is necessary after cold forging and carburizing, and hot forging is essential.
  • Patent Document 4 It is unclear whether the steel for carburized parts of Patent Document 4 can achieve the same effect in cold forging where a large strain is applied, and there is no specific structure form or structure control method. Since it is clear, it cannot be said that the steel shows excellent workability even in the forming forging by applying a large strain in the cold such as plate forging which has been widely applied in recent years.
  • Patent Document 5 does not disclose any knowledge and technique regarding optimum components and microstructures for improving the formability of steel, particularly cold forgeability.
  • the present invention is a steel plate suitable for obtaining parts such as a high cycle gear by sheet forming, and a method for producing the steel plate, which are excellent in cold forgeability and impact resistance after carburizing and quenching and tempering, in view of the above-described prior art.
  • the issue is to provide.
  • the ferrite phase has low hardness and high ductility. Therefore, it is possible to improve the material formability by increasing the grain size in a structure mainly composed of ferrite.
  • carbides in the steel sheet are strong particles that prevent slipping, and by allowing carbides to exist at the ferrite grain boundaries, it is possible to prevent the propagation of slips across the crystal grain boundaries and suppress the formation of shear bands. It can improve the cold forgeability and at the same time improve the formability of the steel sheet.
  • cementite is a hard and brittle structure, and if it exists in the state of pearlite, which is a layered structure with ferrite, the steel becomes hard and brittle, so it must be present in a spherical shape. In consideration of cold forgeability and generation of cracks during forging, the particle size needs to be in an appropriate range.
  • the metal structure of the steel sheet after coiling after hot rolling becomes a bainite structure in which cementite is dispersed in fine pearlite or fine ferrite with a small lamellar spacing, so that the temperature is relatively low (400 ° C to 550 ° C). Take up with.
  • cementite dispersed in the ferrite is also easily spheroidized.
  • the cementite is partially spheroidized by annealing at a temperature just below the Ac1 point as the first stage annealing.
  • annealing is performed at a temperature between Ac1 point and Ac3 point (so-called two-phase region of ferrite and austenite), and a part of the ferrite grains is left, and a part thereof is austenite transformed. Thereafter, the ferrite grains left by slow cooling were grown, and austenite was transformed into ferrite by using the ferrite grains as a nucleus, so that cementite was precipitated at the grain boundaries while obtaining a large ferrite phase, and the above structure was realized.
  • the present invention has been made on the basis of these findings, and the gist thereof is as follows.
  • Component composition is mass%, C: 0.10 to 0.40%, Si: 0.01 to 0.30%, Mn: 0.30 to 1.00%, Al: 0.001 to 0.10%, Cr: 0.50 to 2.00%, Mo: 0.001 to 1.00%, P: 0.020% or less, S: 0.010% or less, N: 0.020% or less O: 0.020% or less, Ti: 0.010% or less, B: 0.0005% or less, Sn: 0.050% or less, Sb: 0.050% or less, As: 0.050% or less, Nb : 0.10% or less, V: 0.10% or less, Cu: 0.10% or less, W: 0.10% or less, Ta: 0.10% or less, Ni: 0.10% or less, Mg: 0 0.050% or less, Ca: 0.050% or less, Y: 0.050% or less, Zr: 0.050% or less, La: 0.050% or less, and Ce: 0 Low carbon steel sheet containing 050%, balance Fe and impurities, where
  • a manufacturing method for manufacturing the steel sheet of (1) wherein the steel slab having the component composition of (1) is hot-rolled to complete finish hot rolling in a temperature range of 650 ° C. to 950 ° C.
  • the hot-rolled steel sheet is wound up at 400 ° C. or higher and 600 ° C. or lower, and the wound hot-rolled steel sheet is pickled, and the pickled hot-rolled steel sheet is 30 ° C./hour or higher and 150 ° C./hour or lower.
  • the first stage annealing is performed for 3 hours or more and 60 hours or less, and then the hot-rolled steel sheet is heated at 1 ° C./hour or more and 80 ° C./hour.
  • the present invention it is possible to provide a steel plate excellent in cold forgeability and impact resistance after carburizing, quenching and tempering, and particularly suitable for obtaining parts such as a high cycle gear by sheet forming and a method for producing the same.
  • % related to the component composition means “mass%”.
  • C is an element that forms carbides in steel and is effective in strengthening steel and refining ferrite grains.
  • it is essential to suppress the coarsening of the ferrite grain size, but if it is less than 0.10%, the volume fraction of carbides Is insufficient, and it becomes impossible to suppress the coarsening of the carbide during annealing, so C is made 0.10% or more. Preferably it is 0.11% or more.
  • the volume fraction of carbide increases, and when a load is instantaneously applied, a large amount of cracks that become the starting point of fracture are generated, leading to a decrease in impact resistance.
  • 0.40% or less Preferably it is 0.38% or less.
  • Si 0.01-0.30%
  • Si is an element that acts as a deoxidizer and affects the morphology of carbides.
  • the austenite phase is generated during annealing by the two-step type annealing, It is necessary to dissolve the carbide and then slowly cool it to promote the formation of carbide at the ferrite grain boundaries.
  • Si exceeds 0.30%, the ductility of ferrite is reduced, cracking is likely to occur during cold forging, and cold forgeability and impact resistance after carburizing and quenching are reduced. % Or less. Preferably it is 0.28% or less.
  • Si is preferably as small as possible, reduction to less than 0.01% causes a significant increase in refining costs, so Si is set to 0.01% or more. Preferably it is 0.02% or more.
  • Mn is an element that controls the form of carbide in the two-step annealing. If it is less than 0.30%, it becomes difficult to form carbides on the ferrite grain boundaries in the slow cooling after the second stage annealing, so Mn is 0.30% or more. Preferably it is 0.33% or more.
  • Mn is made 1.00% or less.
  • Mn is 0.96% or less.
  • Al 0.001 to 0.10%
  • Al is an element that acts as a deoxidizer for steel and stabilizes ferrite. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so Al is made 0.001% or more. Preferably it is 0.004% or more.
  • Al is made 0.10% or less.
  • it is 0.09% or less.
  • Cr and Mo are elements that improve toughness.
  • Cr is an element effective for stabilizing carbide during heat treatment. If it is less than 0.50%, it becomes difficult to leave carbides during carburization, leading to a coarsening of the austenite grain size in the surface layer and a reduction in impact resistance, so Cr is made 0.50% or more. Preferably it is 0.52% or more.
  • the concentration of Cr in the carbide increases, and a lot of fine carbides remain in the austenite phase produced by the two-step annealing, so that the grains after slow cooling Since carbides are also present therein, the increase in hardness and the number ratio of grain boundary carbides are reduced, and the cold forgeability is reduced, so Cr is made 2.00% or less. Preferably it is 1.94% or less.
  • Mo 0.001 to 1.00%
  • Mo is an element effective for controlling the form of carbide. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so Mo is made 0.001% or more. Preferably it is 0.017% or more.
  • Mo is made 1.00% or less. Preferably it is 0.94% or less.
  • the following elements are impurities and must be controlled to a certain amount or less.
  • P 0.020% or less
  • P is an element that segregates at the ferrite grain boundaries and suppresses the formation of grain boundary carbides. The smaller the number, the better.
  • the P content may be 0, but in order to achieve a purity of less than 0.0001% in the refining process, refining takes a long time and causes a significant increase in manufacturing costs. 0.0001 to 0.0013%.
  • P is made 0.020% or less. Preferably it is 0.018% or less.
  • S is an impurity element that forms non-metallic inclusions such as MnS.
  • Non-metallic inclusions are the starting point of cracking during cold forging, so the smaller the S, the better.
  • the content of S may be 0, but if S is reduced to less than 0.0001%, the refining cost increases significantly, so the practical lower limit is 0.0001 to 0.0012%.
  • the crack length during cold forging is increased, so S is made 0.010% or less. Preferably it is 0.009% or less.
  • N 0.020% or less
  • N is an element that segregates to the ferrite grain boundary and suppresses the formation of carbides on the grain boundary. The smaller the number, the better.
  • the N content may be 0, but if it is reduced to less than 0.0001%, the refining cost will increase significantly, so the practical lower limit is 0.0001 to 0.0006%.
  • the ratio of the number of carbides on the ferrite grain boundary to the number of carbides in the ferrite grains is less than 1 even when two-phase annealing and annealing are performed, and the cold forgeability is low. Since N falls, N is made 0.020% or less. Preferably it is 0.017% or less.
  • O is an element that forms an oxide in steel. Since the oxide which exists in a ferrite grain turns into a production
  • the content of O may be 0, but if O is reduced to less than 0.0001%, the refining cost increases significantly, so the practical lower limit is 0.0001 to 0.0006%.
  • the ratio of the number of carbides on the ferrite grain boundary to the number of carbides in the ferrite grains becomes less than 1, and cold forgeability decreases, so O is 0.020% or less. To do. Preferably it is 0.017% or less.
  • Ti 0.010% or less
  • Ti is an element that is important for controlling the form of carbides, and is an element that promotes the formation of carbides in ferrite grains when contained in a large amount.
  • the Ti content may be 0, but if the content is reduced to less than 0.0001%, the refining cost is greatly increased, so the practical lower limit is 0.0001 to 0.0006%.
  • Ti 0.010% or less. To do. Preferably it is 0.007% or less.
  • B is an element effective for controlling dislocation slip during cold forging. Since the activity of the slip system is limited by the inclusion of a large amount, B is preferably as small as possible. The B content may be zero. Careful attention is required for the detection of B less than 0.0001%, and depending on the analyzer, the detection is below the lower limit of detection.
  • B is set to 0.0005% or less. .
  • it is 0.0005% or less.
  • Sn 0.050% or less
  • Sn is an element mixed in from the steel raw material (scrap), and the smaller the better.
  • the Sn content may be 0, but if the content is reduced to less than 0.001%, the refining cost increases significantly, so the practical lower limit is 0.001 to 0.002%.
  • Sn is made 0.050% or less. Preferably it is 0.048% or less.
  • Sb 0.050% or less
  • Sb is an element mixed from steel raw material (scrap) like Sn. Since Sb segregates at the grain boundaries and reduces the number ratio of grain boundary carbides, the smaller the Sb, the better.
  • the Sb content may be 0, but if the content is reduced to less than 0.001%, the refining cost is greatly increased, so the practical lower limit is 0.001 to 0.002%.
  • Sb is made 0.050% or less. Preferably it is 0.048% or less.
  • the content of As may be 0, but if the content is reduced to less than 0.001%, the refining cost increases significantly, so the practical lower limit is 0.001 to 0.002%.
  • the number ratio of grain boundary carbides decreases and cold forgeability decreases, so As is made 0.050% or less.
  • the steel sheet of the present invention uses the above elements as basic elements, but may further contain the following elements for the purpose of improving cold forgeability and other characteristics.
  • the following elements are not essential for obtaining the effects of the present invention, so the content may be zero.
  • Nb is an element effective for controlling the form of carbides, and is an element that contributes to improving toughness by refining the structure. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so Nb is preferably 0.001% or more. More preferably, it is 0.002% or more.
  • Nb is 0 10% or less.
  • V 0.10% or less
  • Nb is an element that is effective in controlling the morphology of carbides, and is an element that contributes to improving toughness by refining the structure. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so V is preferably 0.001% or more. More preferably, it is 0.004% or more.
  • V is 0. 10% or less.
  • Cu 0.10% or less
  • Cu is an element that forms fine precipitates and contributes to improvement in strength. If it is less than 0.001%, the effect of improving the strength cannot be obtained sufficiently, so Cu is preferably made 0.001% or more. More preferably, it is 0.008% or more.
  • Cu is made 0.10% or less. Preferably it is 0.09% or less.
  • W is an element effective for controlling the form of carbide. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so W is preferably 0.001% or more. More preferably, it is 0.003% or more.
  • W is 0. 10% or less. Preferably it is 0.08% or less.
  • Ta 0.10% or less
  • Nb, V, and W is an element effective for controlling the morphology of carbides. If less than 0.001%, the effect of addition cannot be sufficiently obtained, so Ta is preferably made 0.001% or more. Preferably it is 0.007% or more.
  • Ta 0. 10% or less.
  • it is 0.09% or less.
  • Ni 0.10% or less
  • Ni is an element effective for improving the impact resistance characteristics of parts. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so Ni is preferably made 0.001% or more. More preferably, it is 0.002% or more.
  • Ni is made 0.10% or less.
  • it is 0.09% or less.
  • Mg 0.050% or less
  • Mg is an element that can control the form of sulfide by addition of a small amount. If it is less than 0.0001%, the effect of addition cannot be sufficiently obtained, so Mg is preferably made 0.0001% or more. More preferably, it is 0.0008% or more.
  • Mg is made 0.050% or less. Preferably it is 0.049% or less.
  • Ca 0.050% or less
  • Ca is an element that can control the form of sulfide with a small amount of addition. If less than 0.001%, the effect of addition cannot be sufficiently obtained, so Ca is preferably made 0.001% or more. More preferably, it is 0.003% or more.
  • Ca is made 0.050% or less.
  • it is 0.04% or less.
  • Y like Mg and Ca, is an element that can control the form of sulfide by addition of a trace amount. If it is less than 0.001%, the effect of addition cannot be sufficiently obtained, so Y is preferably 0.001% or more. More preferably, it is 0.003% or more.
  • Y is 0.050% or less.
  • it is 0.031% or less.
  • Zr 0.050% or less
  • Zr is an element that can control the form of sulfide by adding a small amount. If it is less than 0.001%, the effect of addition cannot be obtained sufficiently, so Zr is preferably 0.001% or more. More preferably, it is 0.004% or more.
  • Zr is made 0.050% or less.
  • it is 0.045% or less.
  • La is an element that is effective in controlling the form of sulfides when added in a small amount, and is an element that segregates at the grain boundaries and lowers the number ratio of grain boundary carbides. If it is less than 0.001%, the shape control effect cannot be obtained sufficiently, so La is preferably 0.001% or more. More preferably, it is 0.003% or more.
  • La is made 0.050% or less. Preferably it is 0.047% or less.
  • Ce is an element that can control the form of sulfide by addition of a small amount, and is an element that segregates at the grain boundary and lowers the number ratio of grain boundary carbides. If it is less than 0.001%, the shape control effect cannot be obtained sufficiently, so Ce is preferably 0.001% or more. More preferably, it is 0.003% or more.
  • Ce is made 0.050% or less. Preferably it is 0.046% or less.
  • the balance of the component composition of the steel sheet of the present invention is Fe and inevitable impurities.
  • the structure of the steel sheet of the present invention is substantially a structure composed of ferrite and carbide.
  • carbides are compounds in which Fe atoms in cementite are substituted with Mn, Cr, etc., alloy carbides (M 23 C 6 , M 6 C, MC, etc.) , M is Fe and other metal elements).
  • a shear band is formed in the macro structure of the steel sheet, and slip deformation is concentrated in the vicinity of the shear band. Slip deformation is accompanied by dislocation growth, and a region having a high dislocation density is formed in the vicinity of the shear band. As the amount of strain applied to the steel sheet increases, slip deformation is promoted and the dislocation density increases.
  • the formation of a shear band can be understood as a phenomenon in which a slip generated in one grain overcomes the grain boundary and continuously propagates to adjacent grains. Therefore, in order to suppress the formation of shear bands, it is necessary to prevent the propagation of slip across the crystal grain boundary.
  • Carbide in the steel plate is a strong particle that prevents sliding, and by allowing the carbide to be present at the ferrite grain boundary, formation of shear bands can be suppressed and cold forgeability can be improved.
  • the carbide In order to obtain such an effect, the carbide needs to be dispersed in an appropriate size in the metal structure. Therefore, the average particle diameter of the carbide is set to 0.4 ⁇ m or more and 2.0 ⁇ m or less.
  • the particle diameter of the carbide is less than 0.4 ⁇ m, the hardness of the steel sheet is remarkably increased and the cold forgeability is lowered. More preferably, it is 0.6 ⁇ m or more.
  • the average particle diameter of the carbide exceeds 2.0 ⁇ m, the carbide becomes a starting point of cracking during cold forming. More preferably, it is 1.95 ⁇ m or less.
  • cementite which is a carbide of iron
  • the area ratio is set to 6% or less.
  • perlite Since perlite has a unique lamellar structure, it can be distinguished by SEM and optical microscope observation.
  • the area ratio of pearlite can be obtained by calculating the region of the lamellar structure in an arbitrary cross section.
  • cold forgeability is considered to be strongly influenced by the carbide coverage of ferrite grain boundaries, and its high-precision measurement is required, but carbide to ferrite grain boundaries in a three-dimensional space.
  • serial sectioning SEM observation or three-dimensional EBSP observation in which the cutting and observation of the sample by FIB is repeatedly performed in a scanning electron microscope is indispensable. Accumulation is essential.
  • the present inventors have clarified this, and as a result of searching for a simpler and more accurate evaluation index, if the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains is used as an index, the cold The present inventors have found that forgeability can be evaluated, and if the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains exceeds 1, the cold forgeability is significantly improved. .
  • Carbide is observed with a scanning electron microscope. Prior to observation, the tissue observation sample was wet-polished with emery paper and polished with diamond abrasive grains having an average particle size of 1 ⁇ m, and the observation surface was finished to a mirror surface, and then the tissue was saturated with a saturated picric acid-alcohol solution. Etch.
  • the observation magnification is set to 3000 times, and 8 images of a 30 ⁇ m ⁇ 40 ⁇ m visual field in a 1/4 layer thickness are taken at random.
  • the area of each carbide contained in the region is measured in detail with the image analysis software represented by Mitani Corporation (WinWROOF) on the obtained tissue image.
  • carbides having an area of 0.01 ⁇ m 2 or less are excluded from the evaluation target.
  • Cold forgeability can be improved by setting the ferrite grain size to 3.0 ⁇ m or more and 50.0 ⁇ m or less as a structure after annealing. If the ferrite particle size is less than 3 ⁇ m, the hardness increases and cracks and cracks are likely to occur during cold forging, so the ferrite particle size is preferably 3.0 ⁇ m or more. More preferably, it is 7.5 ⁇ m or more.
  • the ferrite grain size is preferably 50.0 ⁇ m or less. . More preferably, it is 37.9 ⁇ m or less.
  • the ferrite grain size was determined by observing the structure of the observation surface etched with a 3% nitric acid-alcohol solution with an optical microscope or scanning electron microscope after polishing the observation surface of the sample for tissue observation to a mirror surface in the above-described procedure.
  • the line segment method is applied to the captured image.
  • the Vickers hardness of the steel sheet By setting the Vickers hardness of the steel sheet to 100 HV or more and 180 HV or less, cold forgeability and impact resistance after carburizing and tempering can be improved. If the Vickers hardness is less than 100 HV, buckling is likely to occur during cold forging, and folding and folding of the buckled portion occur and impact resistance is reduced. Therefore, the Vickers hardness is 100 HV or more. To do. Preferably it is 110HV or more.
  • the Vickers hardness is set to 180 HV or less. Preferably it is 170 HV or less.
  • FIG. 1 schematically shows the outline of the cold forging test and the mode of cracks introduced by cold forging.
  • FIG. 1 (a) shows a disk-shaped test material cut out from a hot-rolled steel sheet
  • FIG. 1 (b) shows the shape of the test material after cold forging
  • FIG. 1 (c) shows the result after cold forging. The cross-sectional aspect of this test material is shown.
  • a disk-shaped test material 1 having a diameter of 70 mm is cut out from a hot-rolled steel sheet having a thickness of 5.2 mm (see FIG. 1 (a)).
  • a test material is prepared (not shown).
  • a cup-shaped test material 2 having a diameter of 30 mm, a height of 30 mm, and a vertical wall thickness of 8 mm is produced (see FIG. 1B).
  • the cup-shaped test material 2 subjected to thickening molding is cut with a FANUC wire-cut electric discharge machine so that the cross-section of the diameter portion appears (see FIG. 1 (c)).
  • the cold forgeability is evaluated by this measured value.
  • the hot-rolled steel plate which this invention makes object is not limited to the hot-rolled steel plate of plate
  • the present invention can improve cold forgeability and impact resistance after carburizing and tempering even in a hot rolled steel sheet having a general thickness (2 to 15 mm).
  • the technical idea of the manufacturing method of the present invention is to consistently manage hot rolling conditions and annealing conditions when manufacturing steel sheets from steel slabs having the above-mentioned composition, and to provide cold forgeability and impact resistance characteristics after carburizing and quenching and tempering. It is to improve.
  • a molten steel having a required composition is continuously cast to form a slab, and the slab is subjected to hot rolling as usual, or after being cooled and heated, and then subjected to hot rolling, 650 ° C. to 950 ° C. Finish hot rolling in the following temperature range.
  • the hot-rolled steel sheet after finish rolling is cooled on the ROT and wound at a winding temperature of 400 ° C. or higher and 600 ° C. or lower.
  • the hot-rolled steel sheet is subjected to two-step annealing that is held in two temperature ranges after pickling. At that time, in the first-stage annealing, the hot-rolled steel sheet is heated to an annealing temperature of 30 ° C./hour or more to 150 ° C./hour or more. Heating is performed at a heating rate of °C / hour or less, and annealing is performed in a temperature range of 650 ° C. to 720 ° C. for 3 hours to 60 hours.
  • the hot-rolled steel sheet is heated to the annealing temperature at a heating rate of 1 ° C./hour to 80 ° C./hour, and in a temperature range of 725 ° C. to 790 ° C. for 3 hours to 50 hours.
  • the following annealing is performed.
  • the annealed hot-rolled steel sheet is cooled to 650 ° C. at a cooling rate of 1 ° C./hour or more and 100 ° C./hour or less, and then cooled to room temperature.
  • the molten steel having the required composition is continuously cast into a slab, as it is, or once cooled and heated, and then subjected to hot rolling, and finish hot rolling is finished at a temperature range of 650 ° C. to 950 ° C.
  • the rolled steel sheet is wound up at 400 ° C. or higher and 600 ° C. or lower.
  • the slab heating temperature is preferably 1300 ° C. or lower, and the heating time for maintaining the slab surface temperature at 1000 ° C. or higher is preferably 7 hours or shorter.
  • the heating temperature is preferably 1300 ° C. or less, and the heating time is preferably 7 hours or less. More preferably, the heating temperature is 1280 ° C. or less, and the heating time is 6 hours or less.
  • Finish hot rolling ends at a temperature of 650 ° C or higher and 950 ° C or lower. If the finishing hot rolling temperature is less than 650 ° C., the rolling load is remarkably increased due to an increase in deformation resistance of the steel material, and further, the roll wear amount is increased and the productivity is lowered. Therefore, the finishing hot rolling temperature is 650 ° C. That's it. Preferably it is 680 degreeC or more.
  • the finish hot rolling temperature exceeds 950 ° C.
  • a thick scale is generated while passing through ROT (Run Out Table), and the surface of the steel sheet is wrinkled due to the scale, and during cold forging, and / or
  • the finish hot rolling temperature is 950 ° C. or lower.
  • it is 920 degrees C or less.
  • the cooling rate when the hot-rolled steel sheet is cooled on the ROT is preferably 10 ° C./second or more and 100 ° C./second or less.
  • the cooling rate is less than 10 ° C./second, the generation of thick scale and the generation of wrinkles due to it cannot be suppressed during the cooling, and the impact resistance is reduced.
  • Seconds or more are preferred. More preferably, it is 20 ° C./second or more.
  • the hot-rolled steel sheet is cooled from the surface layer of the steel sheet to the inside at a cooling rate exceeding 100 ° C./second, the outermost layer part is excessively cooled, and low-temperature transformation structures such as bainite and martensite are formed in the outermost layer part. Arise.
  • the cooling rate is preferably 100 ° C./second or less. More preferably, it is 80 ° C./second or less.
  • the cooling rate is as follows: from the time when the hot-rolled steel sheet after finish hot rolling passes through the non-water-injection section to the time when the water-cooling in the water-injection section is cooled to the winding target temperature on the ROT. It refers to the cooling capacity received from the cooling equipment in the water injection section, and does not indicate the average cooling rate from the water injection start point to the temperature taken up by the winder.
  • the winding temperature is 400 ° C or higher and 600 ° C or lower. This is a temperature lower than the general winding temperature.
  • the structure of the steel sheet can be a bainite structure in which carbides are dispersed in fine ferrite.
  • the coiling temperature is set to 400 ° C. or higher. Preferably it is 430 degreeC or more.
  • the coiling temperature exceeds 600 ° C.
  • pearlite with a large lamellar spacing is generated, and thick needle-like carbides having high thermal stability are formed. Even after the two-step annealing, the needle-like carbides are formed. Remains.
  • the coiling temperature is set to 600 ° C. or lower. Preferably it is 570 degrees C or less.
  • the hot-rolled steel sheet manufactured under the above conditions is subjected to two-step annealing that is held in two temperature ranges after pickling.
  • the stability of the carbide is controlled and the formation of the carbide at the ferrite grain boundary is promoted.
  • the carbide is coarsened and the added gold element is concentrated to enhance the thermal stability of the carbide. Thereafter, the temperature is raised to a temperature range of Ac1 point or higher, austenite is generated in the structure, carbides in fine ferrite grains are dissolved in austenite, and coarse carbides remain in the austenite.
  • the austenite is transformed into ferrite by slow cooling, and the carbon concentration in the austenite is increased.
  • carbon atoms are adsorbed on the carbide remaining in the austenite, and the carbide and austenite cover the ferrite grain boundaries, and finally a structure in which a large amount of carbide exists in the ferrite grain boundaries. Can be formed. Therefore, it is clear that the structure defined in the present invention cannot be formed only by simple annealing.
  • the heating rate up to the first stage annealing temperature is 30 ° C./hour or more and 150 ° C./hour or less. If the heating rate is less than 30 ° C./hour, it takes time to raise the temperature and the productivity decreases, so the heating rate is set to 30 ° C./hour or more. Preferably, it is 40 ° C./hour or more.
  • the heating rate exceeds 150 ° C./hour, the temperature difference between the outer peripheral portion and the inside of the coil increases, so that frost and seizure occur due to the thermal expansion difference, and irregularities are generated on the surface of the steel sheet.
  • the heating rate is set to 150 ° C./hour or less. Preferably it is 120 degrees C / hour or less.
  • the annealing temperature in the first stage annealing is 650 ° C. or more and 720 ° C. or less. If the first stage annealing temperature is less than 650 ° C., the stability of the carbide is insufficient, and it becomes difficult to leave the carbide in the austenite in the second stage annealing, so the first stage annealing temperature is 650 ° C. or higher. Preferably it is 670 degreeC or more.
  • the annealing temperature is set to 720 ° C. or less. Preferably it is 700 degrees C or less.
  • the holding time in the first stage annealing (first stage holding time) is 3 hours or more and 60 hours or less. If the first stage holding time is less than 3 hours, the carbide is not sufficiently stabilized, and it is difficult to leave the carbides in the second stage annealing, so the first stage holding time is 3 hours. That's it. Preferably it is 10 hours or more.
  • the holding time of the first stage exceeds 60 hours, further improvement in the stability of the carbide cannot be expected, and further, the productivity is lowered. Therefore, the holding time of the first stage is set to 60 hours or less. Preferably it is 50 hours or less.
  • the hot-rolled steel sheet After completion of holding in the first stage annealing, the hot-rolled steel sheet is heated to the annealing temperature at a heating rate of 1 ° C./hour to 80 ° C./hour.
  • the ferrite grain size When cooled without performing the second-stage annealing, the ferrite grain size does not increase and an ideal structure cannot be obtained.
  • austenite is generated and grows from the ferrite grain boundary.
  • nucleation of austenite can be suppressed, and the grain boundary coverage of carbides can be increased in the structure obtained after slow cooling. Therefore, it is preferable that the heating rate in the second stage annealing is small.
  • the heating rate is set to 1 ° C / hour or more.
  • it is 10 ° C./hour or more.
  • the heating rate exceeds 80 ° C./hour, the temperature difference between the outer peripheral portion and the inside of the coil increases, resulting in a large difference in thermal expansion due to transformation, causing scouring and seizure, and unevenness on the steel sheet surface. Produces. At the time of cold forging, cracks are generated starting from this unevenness, which causes a decrease in cold forgeability and a decrease in impact resistance after carburizing and quenching and tempering, so the heating rate is 80 ° C./hour or less.
  • the annealing temperature in the second stage annealing is 725 ° C. or higher and 790 ° C. or lower. If the second stage annealing temperature is less than 725 ° C., the amount of austenite produced is reduced, and after cooling after the second stage annealing, the number ratio of carbides on the ferrite grain boundaries is reduced. Becomes smaller. Therefore, the second stage annealing temperature is set to 725 ° C. or higher. Preferably it is 735 ° C or more.
  • the annealing temperature of the second stage exceeds 790 ° C, it becomes difficult to leave the carbide in the austenite, and it becomes difficult to control the above-described structural change, so the annealing temperature of the second stage is 790 ° C or less. And Preferably it is 780 degrees C or less.
  • the holding time in the second stage annealing is 1 hour to 50 hours. If the second stage holding time is less than 1 hour, the amount of austenite produced is small, and the carbides in the ferrite grains are not sufficiently dissolved, and the number ratio of carbides on the ferrite grain boundaries can be increased. Since it becomes difficult and the ferrite grain size becomes small, the second stage holding time is set to 1 hour or more. Preferably it is 5 hours or more.
  • the second stage holding time is set to 50 hours or less. Preferably it is 45 hours or less.
  • Cooling stop temperature 650 ° C
  • Cooling rate 1 ° C / hour or more and 100 ° C / hour or less
  • the annealed hot-rolled steel sheet is gradually cooled to 650 ° C. at a cooling rate of 1 ° C./hour or more and 100 ° C./hour or less.
  • the slow cooling stop temperature exceeds 650 ° C.
  • the untransformed austenite is transformed into pearlite or bainite by the cooling rate exceeding 100 ° C./hour until the room temperature, the hardness increases, and the cold forgeability increases. Since it decreases, the cooling stop temperature is set to 650 ° C.
  • the cooling rate is low.
  • the cooling rate is set to 1 ° C./hour or more.
  • it is 10 ° C./hour or more.
  • the cooling rate exceeds 100 ° C./hour, austenite is transformed into pearlite, the hardness of the steel sheet is increased, and cold forgeability is deteriorated and impact resistance characteristics after carburizing and quenching and tempering are reduced.
  • the cooling rate is 100 ° C./hour or less. Preferably, it is 90 ° C./hour.
  • the cooling stop temperature is a temperature to be controlled at the above cooling rate. If cooling to 650 ° C. is performed at a cooling rate of 1 ° C./hour or more and 100 ° C./hour or less, the temperature is reduced to 650 ° C. or less.
  • the cooling is not particularly limited.
  • the annealing atmosphere is not limited to a specific atmosphere.
  • any of an atmosphere of 95% or more of nitrogen, an atmosphere of 95% or more of hydrogen, and an air atmosphere may be used.
  • the manufacturing method that consistently manages the hot rolling conditions and annealing conditions of the present invention and performs the structure control of the steel sheet, it is excellent in cold forging combined with drawing and thickening forming. It is possible to produce a low carbon steel sheet that exhibits forgeability and also has excellent impact resistance after carburizing, quenching and tempering.
  • a continuous cast slab (steel ingot) having the composition shown in Table 1 was heated at 1240 ° C. for 1.8 hours and then subjected to hot rolling. Finished hot rolling was completed at 890 ° C., cooled to 520 ° C. at a cooling rate of 45 ° C./second on the ROT, wound up at 510 ° C., and a hot-rolled coil having a thickness of 5.2 mm was manufactured.
  • the hot-rolled coil is pickled, charged in a box-type annealing furnace, the atmosphere is controlled to 95% hydrogen-5% nitrogen, and then heated from room temperature to 705 ° C. at a heating rate of 100 ° C./hour.
  • the temperature distribution in the coil was made uniform by maintaining at 705 ° C. for 36 hours.
  • cooling to 650 ° C. at a cooling rate of 10 ° C./hour followed by furnace cooling to room temperature.
  • a sample for characteristic evaluation was prepared.
  • the structure of the sample was observed by the method described above, and the crack length existing in the sample after cold forging was measured by the method described above.
  • Carburization of the thickened sample was performed by gas carburization.
  • a treatment is performed at 940 ° C. for 120 minutes, It was cooled to the furnace.
  • FIG. 2 schematically shows an outline of a drop weight test for evaluating the impact resistance characteristics of a sample subjected to carburizing, quenching and tempering.
  • the bottom of the cup-shaped sample 4 subjected to carburizing, quenching and tempering is fixed with a jig, and a weight of 2 kg is dropped on the cup side (upper side width: 50 mm, lower side width: 10 mm, height: 80 mm, length: 110 mm). Is dropped from the upper part 4 m away, an impact of about 80 J is applied to the vertical wall of the sample 4, the presence or absence of cracking of the sample is observed, and the impact resistance characteristics are evaluated.
  • Table 2 shows the carbide diameter, pearlite area ratio, ferrite grain size, Vickers hardness, ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains, and the maximum thickness of the vertical wall portion in the manufactured sample. The ratio of crack length and the measurement results and evaluation results of impact resistance are shown.
  • the invention steels A-1, B-1, C-1, D-1, E-1, F-1, G-1, H-1, I-1, J-1, and , K-1 has a ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains exceeding 1, Vickers hardness is 100HV or more and 180HV or less, cold forgeability and carburizing and quenching and tempering. Excellent impact resistance after.
  • Comparative Steel L-1 has a low C content and a hardness before cold forging of less than 100 HV, so the cold forgeability is low.
  • Comparative steels M-1, P-1, and Z-1 contain excessive amounts of P, Al, and N, and have a large amount of segregation at the ⁇ / ⁇ interface during the second stage annealing. The formation of carbide is suppressed.
  • Comparative Steel S-1 contains Si excessively and has low ductility of ferrite, so that cold forgeability is low. Since the comparative steels N-1 and T-1 each contain excessive Mo and Cr, carbides are finely dispersed in the ferrite grains, and the hardness exceeds 180 HV. Since the comparative steel Q-1 contains an excessive amount of Mn, the impact resistance after carburizing and tempering is remarkably low.
  • Comparative steel O-1 has a low amount of Cr, and the austenite grains on the surface layer are abnormally coarsened during carburizing, so the impact resistance is low. Since the comparative steel R-1 contains excessive S, coarse MnS is generated in the steel and the cold forgeability is low. Since comparative steel U-1 contains C excessively, coarse carbides are generated inside the thickened steel, and the coarse carbides remain after carburizing and quenching, so that the impact resistance is low.
  • Comparative Steel V-1 has a low Mn content and it was difficult to increase the stability of carbides, and therefore, cold forgeability and impact resistance after carburizing and tempering were low. Since the comparative steels W-1 and X-1 contain excessive amounts of O and Ti, the oxides and TiC present in the ferrite grains become carbide generation sites in the slow cooling after the two-phase annealing, and the grain boundaries The formation of carbides in is suppressed, and the cold forgeability is low. Since comparative steel Y-1 contains B excessively, cold forgeability is low.
  • Table 4 shows the carbide diameter, pearlite area ratio, ferrite particle diameter, Vickers hardness, ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains, and the thickness of the vertical wall portion of the prepared sample. The ratio of the maximum crack length and the measurement results and evaluation results of impact resistance are shown.
  • Comparative steel E-3 has a low finish hot rolling temperature, an increased rolling load, and low productivity.
  • Comparative Steel D-2 has a high finish hot rolling temperature and scale flaws formed on the surface of the steel sheet. Therefore, when subjected to the wear resistance test after quenching and tempering, cracks and peeling occurred starting from the scale flaw, resulting in wear resistance. The characteristics deteriorated.
  • the comparative steel F-2 has a slow cooling rate in ROT (Run Out Table), which causes a decrease in productivity and generation of scale defects.
  • Comparative Steel C-4 the cooling rate at the ROT was 100 ° C./second, and the outermost layer portion of the steel sheet was excessively cooled, so that fine cracks were generated in the outermost layer portion.
  • the comparative steel C-2 has a low coiling temperature, a large amount of low-temperature transformation structures such as bainite and martensite are generated and become brittle, cracks occur frequently when the hot-rolled coil is discharged, and productivity is lowered. Furthermore, the wear resistance characteristics of samples taken from the cracks are low.
  • the comparative steel G-2 has a high coiling temperature, and pearlite with thick lamellar spacing is generated in the hot rolled structure, and the thermal stability of the needle-like coarse carbide is high, and after the two-step type annealing, However, since the carbides remain in the steel sheet, the machinability is low.
  • the comparative steel H-4 has low productivity because the heating rate in the first stage annealing of the two-step type annealing is slow.
  • comparative steel E-3 Since comparative steel E-3 has a high heating rate in the first stage annealing, the temperature difference between the inside of the coil and the inner and outer peripheral portions becomes large, and sore and seizure due to the difference in thermal expansion occurs and quenching occurs. When subjected to an evaluation test of wear resistance after tempering, cracks and peeling occurred from the heel part, and the wear resistance was lowered.
  • Comparative Steel G-4 has a low holding temperature (annealing temperature) in the first stage annealing, the carbide coarsening treatment below the Ac1 point is insufficient, and the thermal stability of the carbide is insufficient. Thus, the remaining carbides in the second stage annealing are reduced, and the pearlite transformation cannot be suppressed in the structure after the slow cooling, so that the machinability is low.
  • annealing temperature annealing temperature
  • Comparative Steel D-4 has a high holding temperature (annealing temperature) in the first stage annealing, austenite is generated during annealing, and the stability of the carbide cannot be increased. Therefore, pearlite is generated after annealing, and Vickers hard Is over 180 HV, and machinability is low. Comparative steel J-4 has a short holding time in the first-stage annealing, cannot improve the stability of the carbide, and has low machinability.
  • the comparative steel F-2 has a long holding time in the first stage annealing, low productivity, seizure flaws, and low wear resistance.
  • the comparative steel B-4 has a low productivity because the heating rate in the second stage annealing of the two-step type annealing is slow.
  • Comparative Steel A-3 has a high heating rate in the second-stage annealing, so the temperature difference between the inside and the outer periphery of the coil is large, and the soot and seizure due to a large difference in thermal expansion due to transformation occurs. Low wear resistance after quenching and tempering.
  • Comparative steel K-2 has a low holding temperature (annealing temperature) in the second stage annealing, a small amount of austenite is generated, and the number ratio of carbides at the ferrite grain boundaries cannot be increased, so that the machinability is low.
  • Comparative Steel C-4 has a high holding temperature (annealing temperature) in the second-stage annealing, and the dissolution of carbides during the annealing is promoted, making it difficult to form grain boundary carbides after slow cooling. Produced, Vickers hardness exceeds 180HV, and machinability is low.
  • Comparative steel J-3 has a long holding time in the second stage annealing and promotes dissolution of carbides, and therefore has low machinability.
  • the cooling rate from the second stage annealing to 650 ° C. is slow, the productivity is low, and coarse carbides are generated in the structure after the slow cooling. Cracks occurred as a starting point, and cold forgeability deteriorated.
  • the comparative steel I-3 has a high cooling rate from the second annealing to 650 ° C., and pearlite transformation occurs during cooling to increase the hardness, so that the cold forgeability is low.
  • a continuous cast slab (steel ingot) having the composition shown in Table 5 and Table 6 (continuation of Table 5) was heated at 1240 ° C. for 1.8 hours, It was subjected to hot rolling. Finished hot rolling was completed at 890 ° C., cooled to 520 ° C. at a cooling rate of 45 ° C./second on the ROT, wound up at 510 ° C., and a hot-rolled coil having a thickness of 5.2 mm was manufactured.
  • the temperature from room temperature to 705 ° C. was increased at a heating rate of 100 ° C./hour.
  • Heat and hold at 705 ° C. for 36 hours to homogenize the temperature distribution in the coil then heat to 760 ° C. at a heating rate of 5 ° C./hour, further hold at 760 ° C. for 10 hours, and then to 650 ° C.
  • the structure of the sample was observed by the above method, and the crack length existing in the sample after cold forging was measured by the above method.
  • Table 7 shows the carbide diameter, pearlite area ratio, ferrite grain size, Vickers hardness, ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grain, and the maximum thickness of the vertical wall portion in the manufactured sample. The ratio of crack length and the measurement results and evaluation results of impact resistance are shown.
  • invention steels AA-1, AB-1, AC-1, AD-1, AE-1, AF-1, AG-1, AH-1, AI-1, AJ-1, AK -1, AL-1, AM-1, AN-1, AO-1, AP-1, and AQ-1 are all the ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains Is more than 1, Vickers hardness is 100HV or more and 180HV or less, and is excellent in cold forgeability and impact resistance after carburizing and tempering.
  • the comparative steels AR-1, AS-1, AW-1, AZ-1, BB-1, and BF-1 each contain excessive amounts of La, As, Cu, Ni, Sb, and Ce.
  • the comparative steel BG-1 contains an excessive amount of Si and has a low ductility of ferrite, so that the cold forgeability is low.
  • Comparative steels AT-1, AV-1, BA-1, BC-1, BH-1, and BJ-1 each contain excessive amounts of Mo, Nb, Cr, Ta, W, and V, so ferrite
  • the carbide is finely dispersed in the grains, and the hardness exceeds 180 HV. Since the comparative steel BF-1 contains excessive Mn, the impact resistance after carburizing and tempering is remarkably low.
  • the comparative steels AU-1, AX-1, AY-1, and BE-1 contain excessive amounts of Zr, Ca, Mg, and Y, respectively, and coarse oxides or non-metallic inclusions are formed in the steel. Then, cracks occurred starting from coarse oxides or coarse non-metallic inclusions during cold forging, and cold forgeability deteriorated.
  • the comparative steel BD-1 contains excessive Sn, the ferrite becomes brittle, and the cold forgeability is low. Since the comparative steel BK-1 contains an excessive amount of C, coarse carbides are generated inside the thickened steel, and the coarse carbides remain after carburizing and quenching, resulting in a reduction in impact resistance.
  • a hot-rolled sheet annealing sample having a sheet thickness of 5.2 mm was produced from the slab having the composition of the following conditions under the hot-rolling conditions and annealing conditions shown in Table 8.
  • Table 9 shows the carbide diameter, pearlite area ratio, ferrite particle diameter, Vickers hardness, ratio of the number of carbides in the ferrite grain boundary to the number of carbides in the ferrite grains, and the maximum thickness of the vertical wall portion in the prepared sample. The ratio of crack length and the measurement results and evaluation results of impact resistance are shown.
  • Comparative steel AC-2 has a low finish hot rolling temperature and low productivity.
  • the comparative steel AN-4 has a high finish hot rolling temperature, scale flaws formed on the surface of the steel sheet, and when an impact load is applied after cold forging and carburizing quenching and tempering, cracks are generated from the flaws and impact resistance The characteristics deteriorated.
  • Invented steel AB-3 has a slow cooling rate in the ROT, resulting in a decrease in productivity and a derivation of scale defects.
  • Inventive steels AJ-3 and AD-4 had a cooling rate at ROT of 100 ° C./second, and the outermost layer portion of the steel sheet was excessively cooled, and fine cracks were generated in the outermost layer portion.
  • the comparative steel AN-3 had a low coiling temperature, a lot of low-temperature transformation structures such as bainite and martensite were formed and became brittle, and cracks occurred frequently when the hot-rolled coil was discharged, resulting in decreased productivity. Furthermore, the impact resistance characteristics after cold forging and carburizing quenching and tempering in samples taken from the cracked pieces were inferior.
  • the comparative steel AH-3 has a high coiling temperature, and a pearlite with a thick lamellar spacing is generated in the hot rolled structure, and the thermal stability of the needle-like coarse carbide is high, even after annealing of the two-step type. Since the carbide remains in the steel sheet, the cold forgeability is low.
  • Comparative steel AF-4 has low productivity because the heating rate in the first stage annealing of the two-step type annealing is slow. Since the comparative steel AG-2 has a high heating rate in the first stage annealing, the temperature difference between the inside and the outer periphery of the coil becomes large, so that there is a crack and seizure due to the difference in thermal expansion. Impact resistance after forging and carburizing quenching and tempering decreased.
  • the holding temperature (annealing temperature) in the first stage annealing is low, the coarsening treatment of the carbide below the Ac1 point is insufficient, and the thermal stability of the carbide is insufficient.
  • the carbide remaining at the time of annealing of the eyes decreased, the pearlite transformation could not be suppressed in the structure after the slow cooling, and the cold forgeability was lowered.
  • the comparative steel AM-3 has a high holding temperature (annealing temperature) in the first stage, austenite is generated during annealing, the stability of the carbide cannot be increased, cold forgeability and resistance after carburizing and quenching and tempering. Impact characteristics deteriorated.
  • Comparative steel AF-2 has a short holding time in the first-stage annealing, cannot improve the stability of the carbide, and has low cold forgeability.
  • the comparative steel AO-4 has a long holding time in the first stage annealing and low productivity.
  • Comparative steel AP-4 has low productivity because the heating rate in the second stage annealing of the two-step type annealing is slow. Since the comparative steel AI-3 has a high heating rate in the second stage of annealing, the temperature difference between the inside and the outer periphery of the coil is large, and there is a crack and seizure due to a large difference in thermal expansion due to transformation. When an impact load was applied after quenching and tempering, cracks were generated from the ridges, and the impact resistance characteristics deteriorated.
  • the comparative steel AL-3 has a low holding temperature (annealing temperature) in the second stage annealing, produces a small amount of austenite, cannot increase the number of carbides at the ferrite grain boundaries, and has a cold forgeability. Declined. Since the comparative steel AD-2 has a high holding temperature (annealing temperature) in the second stage annealing, and the dissolution of carbides is accelerated during annealing, it becomes difficult to produce grain boundary carbides after slow cooling, and cold forgeability And the impact resistance after carburizing quenching and tempering decreased.
  • annealing temperature annealing temperature
  • Comparative steel AJ-4 has a long holding time in the second-stage annealing, and promotes the dissolution of carbides, so that the cold forgeability is low.
  • the cooling rate from the second stage annealing to 650 ° C. is slow, the productivity is low, and coarse carbides are generated in the structure after the slow cooling.
  • the comparative steel AP-2 had a high cooling rate from the second stage annealing to 650 ° C., pearlite transformation occurred during cooling, and the hardness increased, resulting in a decrease in cold forgeability.
  • FIG. 3 shows the relationship between the ratio of the number of intergranular carbides to the number of intragranular carbides, the crack length of the cold forged specimen and the impact resistance after carburizing and tempering.
  • FIG. 4 shows another relationship between the ratio of the number of intergranular carbides to the number of intragranular carbides, the crack length of the cold forged specimen, and the impact resistance after carburizing and tempering.
  • FIG. 4 is a diagram showing that the crack length can be suppressed even in a steel sheet to which an additive element is added.
  • the present invention it is possible to provide a low carbon steel sheet excellent in cold forgeability and impact resistance after carburizing and tempering, and a method for producing the same. Since the steel sheet of the present invention is suitable as a material for obtaining parts such as high cycle gears by forming by cold forging such as sheet forming, the present invention has high industrial applicability. is there.

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Abstract

L'invention concerne une tôle d'acier à faible teneur en carbone qui présente une excellente forgeabilité à froid et d'excellentes caractéristiques de résistance aux chocs après cémentation au carbone et trempe/revenu, et qui est caractérisée en ce qu'elle a une composition constituante prescrite, présentant un diamètre de particule de carbure moyen allant de 0,4 à 2,0 µm, ayant un rapport de surface de perlite inférieur ou égal à 6 %, ayant un rapport du nombre de carbures au niveau des limites des grains de ferrite au nombre de carbures à l'intérieur des grains de ferrite supérieur à 1, et ayant une dureté Vickers allant de 100 à 180 HV.
PCT/JP2016/065509 2015-05-26 2016-05-25 Tôle d'acier et son procédé de production WO2016190370A1 (fr)

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EP16800074.3A EP3305929A4 (fr) 2015-05-26 2016-05-25 Tôle d'acier et son procédé de production
US15/576,177 US20180230582A1 (en) 2015-05-26 2016-05-25 Steel plate and method of production of same
BR112017024692-9A BR112017024692A2 (pt) 2015-05-26 2016-05-25 placa de aço e método de produção da mesma
MX2017015016A MX2017015016A (es) 2015-05-26 2016-05-25 Placa de acero y metodo de produccion de la misma.
JP2016559466A JP6119923B1 (ja) 2015-05-26 2016-05-25 鋼板及びその製造方法
CN201680030099.7A CN107614727B (zh) 2015-05-26 2016-05-25 钢板及其制造方法
KR1020177033291A KR102029565B1 (ko) 2015-05-26 2016-05-25 강판 및 그의 제조 방법

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US20180171445A1 (en) * 2015-06-17 2018-06-21 Nippon Steel & Sumitomo Metal Corporation Steel plate and method of production of same
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JP6927427B2 (ja) * 2019-01-30 2021-08-25 Ｊｆｅスチール株式会社 高炭素熱延鋼板およびその製造方法
WO2021260954A1 (fr) * 2020-06-26 2021-12-30 日本製鉄株式会社 Matériau d'acier et pièce en acier cémenté
RU2758716C1 (ru) * 2020-08-20 2021-11-01 Публичное акционерное общество «Северсталь» (ПАО "Северсталь") Способ производства горячекатаного проката из инструментальной стали
RU2765047C1 (ru) * 2020-12-28 2022-01-25 Публичное акционерное общество «Северсталь» (ПАО «Северсталь») Способ производства листов толщиной 2-20 мм из высокопрочной износостойкой стали (варианты)
CN115612924B (zh) * 2022-09-19 2023-09-12 攀钢集团攀枝花钢铁研究院有限公司 一种铅铋堆用铁素体/马氏体耐热钢及其制备方法

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KR20190109463A (ko) * 2017-02-21 2019-09-25 제이에프이 스틸 가부시키가이샤 고탄소 열연 강판 및 그 제조 방법
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JPWO2019044970A1 (ja) * 2017-08-31 2019-11-07 日本製鉄株式会社 浸炭用鋼板、及び、浸炭用鋼板の製造方法
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WO2019044971A1 (fr) * 2017-08-31 2019-03-07 新日鐵住金株式会社 Tôle d'acier pour cémentation et procédé de production de tôle d'acier pour cémentation
WO2019044970A1 (fr) 2017-08-31 2019-03-07 新日鐵住金株式会社 Tôle d'acier pour carburation et procédé de fabrication de tôle d'acier pour carburation
US11639536B2 (en) 2017-08-31 2023-05-02 Nippon Steel Corporation Steel sheet for carburizing, and method for manufacturing steel sheet for carburizing
KR102219032B1 (ko) 2017-08-31 2021-02-23 닛폰세이테츠 가부시키가이샤 침탄용 강판, 및 침탄용 강판의 제조 방법
US10934609B2 (en) 2017-08-31 2021-03-02 Nippon Steel Corporation Steel sheet for carburizing, and method for manufacturing steel sheet for carburizing
KR102235355B1 (ko) 2017-08-31 2021-04-02 닛폰세이테츠 가부시키가이샤 침탄용 강판, 및 침탄용 강판의 제조 방법
CN108251756B (zh) * 2017-12-04 2019-01-29 广东精铟海洋工程股份有限公司 一种铌微合金化低温高性能钢及其制备方法
CN108251756A (zh) * 2017-12-04 2018-07-06 广东精铟海洋工程股份有限公司 一种铌微合金化低温高性能钢及其制备方法
JP2020002447A (ja) * 2018-06-29 2020-01-09 Ｊｆｅスチール株式会社 浸炭部材
JP6587038B1 (ja) * 2018-10-02 2019-10-09 日本製鉄株式会社 浸炭用鋼板、及び、浸炭用鋼板の製造方法
KR20200039611A (ko) 2018-10-02 2020-04-16 닛폰세이테츠 가부시키가이샤 침탄용 강판, 및 침탄용 강판의 제조 방법
WO2020070810A1 (fr) * 2018-10-02 2020-04-09 日本製鉄株式会社 Tôle d'acier pour cémentation et procédé de production de tôle d'acier pour cémentation

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CN107614727B (zh) 2020-01-14
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JP6119923B1 (ja) 2017-04-26
TWI604071B (zh) 2017-11-01
US20180230582A1 (en) 2018-08-16
KR20170138508A (ko) 2017-12-15
JPWO2016190370A1 (ja) 2017-06-15
TW201708569A (zh) 2017-03-01
MX2017015016A (es) 2018-04-13
KR102029565B1 (ko) 2019-10-07
CN107614727A (zh) 2018-01-19
EP3305929A9 (fr) 2019-07-17

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