EP3305929A1 - Stahlblech und verfahren zur herstellung davon - Google Patents

Stahlblech und verfahren zur herstellung davon Download PDF

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Publication number
EP3305929A1
EP3305929A1 EP16800074.3A EP16800074A EP3305929A1 EP 3305929 A1 EP3305929 A1 EP 3305929A1 EP 16800074 A EP16800074 A EP 16800074A EP 3305929 A1 EP3305929 A1 EP 3305929A1
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steel
carbides
steel plate
annealing
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EP16800074.3A
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English (en)
French (fr)
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EP3305929A4 (de
EP3305929A9 (de
Inventor
Kengo Takeda
Kazuo HIKIDA
Ken Takata
Motonori Hashimoto
Toshimasa Tomokiyo
Yasushi Tsukano
Takashi Aramaki
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Nippon Steel Corp
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Nippon Steel and Sumitomo Metal Corp
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Publication of EP3305929A1 publication Critical patent/EP3305929A1/de
Publication of EP3305929A4 publication Critical patent/EP3305929A4/de
Publication of EP3305929A9 publication Critical patent/EP3305929A9/de
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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
    • 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
    • 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
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—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
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
    • C—CHEMISTRY; METALLURGY
    • 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
    • C22C38/44—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 steel plate and a method of production of the same.
  • Steel plate containing, by mass%, carbon in an amount of 0.1 to 0.4% is being used as a material for gears, clutches, and other drive system parts of automobiles by being used press-formed, enlarging holes, bent, drawn, thickened, and thinned and cold forged by combinations of the same from a blank.
  • press-formed, enlarging holes, bent, drawn, thickened, and thinned and cold forged by combinations of the same from a blank Compared with conventional hot forging etc., with cold forging, there is the problem that the amount of strain accumulated in the material becomes higher, cracks of the material and buckling at the time of shaping are invited, and deterioration of the part characteristics is caused.
  • PLT 1 discloses, as steel for machine structural use improving toughness by suppressing coarsening of crystal grains in carburization heat treatment, steel for machine structural use containing, by mass%, C: 0.10 to 0.30%, Si: 0.05 to 2.0%, Mn: 0.10 to 0.50%, P: 0.030% or less, S: 0.030% or less, Cr: 1.80 to 3.00%, Al: 0.005 to 0.050%, Nb: 0.02 to 0.10%, and N: 0.0300% or less and having a balance of Fe and unavoidable impurities, having a structure before cold working comprised of ferrite and pearlite structures, and having an average grain size of ferrite grains of 15 ⁇ m or more.
  • PLT 2 discloses, as steel excellent in cold workability and carburizing and quenching ability, steel containing 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, and B: 0.005 to 0.050%, having a balance of Fe and unavoidable impurities, and having a structure mainly comprised of ferrite phases and graphite phases.
  • PLT 3 discloses a steel material for carburized bevel gear use excellent in impact strength, a high toughness carburized bevel gear, and a method of production of the same.
  • PLT 4 discloses steel for carburized part use having excellent workability while suppressing coarsening of crystal grains even with subsequent carburization and having an excellent impact resistance characteristic and impact fatigue resistance characteristic in a part produced by spheroidal annealing, then a cold forging and a carburizing, quenching, and tempering process.
  • PLT 5 discloses as cold tool steel for plasma carburization use a steel containing 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%, further containing one or more of Ni: 0.10 to 2.50%, Cr: 0.1 to 2.0%, and Mo: 3.0% or less, and having a balance of Fe and unavoidable impurities.
  • PLT 1 Japanese Patent Publication No. 2013-040376A
  • the structure of the steel for machine structural use of PLT 1 is a structure of ferrite + pearlite.
  • This structure compared with a ferrite + cementite structure, has a large hardness, so wear of the die in cold forging cannot be suppressed and the steel cannot necessarily be said to be steel for machine structural use excellent in cold forgeability.
  • PLT 3 The method of production of PLT 3 requires further hot forging after cold forging and carburizing. Since hot forging is essential, this is not a method of production leading to fundamentally lower costs.
  • PLT 5 does not disclose at all the findings and art relating to the optimum components and form of structure for improving the formability of steel, in particular cold forgeability.
  • the present invention in consideration of the above prior art, has as its problem the provision of steel plate excellent in cold forgeability and impact resistance characteristic after carburizing, quenching, and tempering, in particular suitable for obtaining a high cycle gear or other part by forming a plate and of a method of production of the same.
  • a ferrite phase is low in hardness and high in ductility. Therefore, in a structure mainly comprised of ferrite, it becomes possible to increase the grain size so as to raise the formability of the material.
  • Carbides by being made to suitably disperse in the metal structure, can maintain the formability of the material while imparting an excellent wear resistance and rolling fatigue characteristic, so provides a structure essential for drive system parts. Further, the carbides in the steel plate are strong particles obstructing slip. By forming carbides at the ferrite grain boundaries, it is possible to prevent propagation of slip exceeding the crystal grain boundaries and suppress the formation of shear zones. Thus the cold forgeability is improved and, simultaneously, the formability of steel plate is also improved.
  • cementite is a hard, brittle structure. If a laminar structure with ferrite present, that is, in the state of pearlite, the steel becomes hard and brittle, so it has to be present in a spheroidal form. If considering the cold forgeability and the occurrence of fractures at the time of forging, its grain size has to be a suitable range.
  • the steel plate is coiled at a relatively low temperature (400°C to 550°C). By coiling at a relatively low temperature, the cementite dispersed in the ferrite also easily becomes spheroidal.
  • the cementite is partially made spheroidal by annealing at a temperature just under the Ac1 point as first stage annealing.
  • part of the ferrite grains is left while part is transformed to austenite by annealing at a temperature between the Ac1 point and Ac3 point (so-called dual phase region of ferrite and austenite).
  • dual phase region of ferrite and austenite By then making the remaining ferrite grains grow while slowly cooling the steel while using these as nuclei to transform the austenite to ferrite, it is possible to obtain large ferrite phases and make cementite precipitate at the grain boundaries to realize the above structure.
  • the present invention was made based on these discoveries and has as its gist the following:
  • a method of production of the steel plate according to (1) comprising the steps of: hot rolling a steel slab of a chemical composition according to claim 1, completing finish hot rolling in a 650°C to 950°C temperature region to obtain a hot rolled steel plate; coiling the hot rolled steel plate at 400°C to 600°C; pickling the coiled hot rolled steel plate and heating the pickled hot rolled steel plate by a 30°C/hour to 150°C/hour heating rate to a 650°C to 720°C annealing temperature and holding it there for 3 hours to 60 hours as first stage annealing; then heating the hot rolled steel plate to an annealing temperature of 725°C to 790°C by a heating rate of 1°C/hour to 80°C/hour and holding the steel plate for 3 hours to 50 hours as second stage annealing; and cooling the annealed hot rolled steel plate to 650°C by a cooling rate of 1°C/hour to 100°C/hour.
  • a steel plate excellent in cold forgeability and impact resistance characteristic after carburizing, quenching, and tempering in particular one suitable for obtaining a high cycle gear or other part by forming a plate.
  • C is an element forming carbides in steel and effective for strengthening the steel and refining the ferrite grains.
  • suppression of coarsening of the ferrite grain size is essential, but if less than 0.10%, the carbides become insufficient in volume fraction and coarsening of the carbides during annealing can no longer be suppressed, so C is made 0.10% or more. Preferably it is 0.11% or more.
  • C is made 0.40% or less.
  • Si is an element which acts as a deoxidizing agent and further has an effect on the form of the carbides.
  • Si exceeds 0.30%, the ferrite falls in ductility, fractures are easily formed at the time of cold forging, and the cold forgeability and impact resistance characteristic after carburizing, quenching, and tempering deteriorate, so Si is made 0.30% or less. Preferably it is 0.28% or less.
  • Si is preferably as low as possible, but reduction to less than 0.01% invites a large increase in refining costs, so Si is made 0.01% or more. Preferably it is 0.02% or more.
  • Mn is an element controlling the form of carbides in two-stage step type annealing. If less than 0.30%, in the gradual cooling after second stage annealing, it becomes difficult to form carbides at the ferrite grain boundaries, so Mn is made 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 is an element acting as a deoxidizing agent of steel and stabilizing ferrite. If less than 0.001%, the effect of addition is not sufficiently obtained, so Al is made 0.001% or more. Preferably it is 0.004% or more.
  • Al is made 0.10% or less. Preferably it is 0.09% or less.
  • Cr and Mo are elements which improve the toughness.
  • Cr is an element effective for stabilization of carbides at the time of heat treatment. If less than 0.50%, it becomes difficult to cause carbides to remain at the time of carburization, coarsening of the austenite grain size at the surface layer is invited, and a drop in the impact resistance characteristic is caused, so Cr is made 0.50% or more. Preferably it is 0.52% or more.
  • Mo is an element effective for control of the form of carbides. If less than 0.001%, the effect of addition is not sufficiently obtained, so Mo is made 0.001% or more. Preferably it is 0.017% or more.
  • Mo concentrates in the carbides and stable carbides increase in the austenite phase as well, so after gradual cooling, carbides are present in the grains as well, an increase in hardness and drop in number ratio of carbides at the grain boundaries are invited, and the cold forgeability falls, so Mo is made 1.00% or less. Preferably it is 0.94% or less.
  • the following elements are impurities and have to be controlled to certain amounts or less.
  • P is an element segregating at the ferrite grain boundaries and suppressing the formation of carbides at the grain boundaries.
  • the smaller amount is preferable.
  • the content of P may also be 0, but a long time is required for refining in order to make the purity a high one of less than 0.0001% in a refining process and a large increase in the manufacturing cost is invited, so the de facto lower limit is 0.0001 to 0.0013%.
  • P is made 0.020% or less. Preferably it is 0.018% or less.
  • S is an impurity element forming MnS and other nonmetallic inclusions.
  • the nonmetallic inclusions form starting points of formation of fractures at the time of cold forging, so the smaller the S, the better.
  • the content of S may also be 0, but to lower S to less than 0.0001%, the refining costs greatly increase, so the de facto lower limit is 0.0001 to 0.0012%.
  • N is an element segregating at the ferrite grain boundaries and suppressing the formation of carbides at the grain boundaries.
  • the smaller amount is preferable.
  • the content of N may also be 0, but if reducing it to less than 0.0001%, the refining costs greatly increase, so the de facto lower limit is 0.0001 to 0.0006%.
  • N is made 0.020% or less. Preferably it is 0.017% or less.
  • O is an element forming oxides in the steel.
  • the oxides present in the ferrite grains become sites for production of carbides, so the smaller the amount, the better.
  • the content of O may also be 0, but if reducing O to less than 0.0001%, the refining costs greatly increase, so the de facto lower limit is 0.0001 to 0.0006%.
  • the ratio of the number of carbides at the ferrite grain boundaries with respect to the number of carbides in the ferrite grains becomes less than 1 and the cold forgeability falls, so O is made 0.020% or less. Preferably it is 0.017% or less.
  • Ti is an element important for control of the form of the carbides. It is an element by which, by inclusion in a large amount, formation of carbides in the ferrite grains is promoted. The smaller amount is preferable.
  • the content of Ti may also be 0, but if reducing it to less than 0.0001%, the refining costs greatly increase, so the de facto lower limit is 0.0001 to 0.0006%.
  • Ti is made 0.010% or less. Preferably it is 0.007% or less.
  • B is an element effective for control of slip of dislocations at the time of cold forging. By inclusion of a large amount, activity of the slip system is limited, so the smaller the amount of B, the better.
  • the content of B may also be 0. Fine care is required for detection of less than 0.0001% of B. Depending on the analysis device, it is below the lower limit of detection.
  • B is made 0.0005% or less. Preferably it is 0.0005% or less.
  • Sn is an element entering from the steel starting materials (scraps). The smaller amount is preferable.
  • the content of Sn may also be 0, but if reducing it to less than 0.001%, the refining costs greatly increase, so the de facto lower limit is 0.001 to 0.002%.
  • Sn is made 0.050% or less. Preferably, it is 0.048% or less.
  • Sb like Sn, is an element entering from the steel starting materials (scraps). Sb segregates at the grain boundaries and lowers the number ratio of carbides at the grain boundaries, so the smaller the amount, the better.
  • the content of Sb may also be 0, but if reducing it to less than 0.001%, the refining costs greatly increase, so the de facto 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 also be 0, but if reducing it to less than 0.001%, the refining cost greatly increases, so the de facto lower limit is 0.001 to 0.002%.
  • the number ratio of carbides at the grain boundaries falls and the cold forgeability falls, so As is made 0.050% or less. Preferably it is 0.045% or less.
  • the steel plate of the present invention has the above elements as basic elements, but may further contain the following elements for the purpose of improving the cold forgeability and other characteristics.
  • the following elements are not essential for obtaining the effects of the present invention, so the contents may also be 0.
  • Nb is an element effective for control of the form of the carbides. Further, it is an element refining the structure and contributing to improvement of the toughness. If less than 0.001%, the effect of addition is not sufficiently obtained, so Nb is preferably made 0.001% or more. More preferably, it is 0.002% or more.
  • Nb is made 0.10% or less. Preferably it is 0.09% or less.
  • V 0.10% or less
  • V is an element effective for control of the form of the carbides. Further, it is an element refining the structure and contributing to improvement of the toughness. If less than 0.001%, the effect of addition is not sufficiently obtained, so V is preferably made 0.001% or more. More preferably, it is 0.004% or more.
  • V is made 0.10% or less.
  • it is 0.09% or less.
  • Cu is an element forming fine precipitates and contributing to improvement of the strength. If less than 0.001%, the effect of improvement of the strength is not sufficiently obtained, 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 control of the form of the carbides. If less than 0.001%, the effect of addition is not sufficiently obtained, so W is preferably made 0.001% or more. More preferably, it is 0.003% or more.
  • W is made 0.10% or less. Preferably, it is 0.08% or less.
  • Ta 0.10% or less
  • Ta is an element effective for control of the form of the carbides. If less than 0.001%, the effect of addition is not sufficiently obtained, so Ta is preferably made 0.001% or more. More preferably, it is 0.007% or more.
  • Ta is made 0.10% or less. Preferably, it is 0.09% or less.
  • Ni is an element effective for improvement of the impact resistance characteristic of parts. If less than 0.001%, the effect of addition is not sufficiently obtained, so Ni preferably is made 0.001% or more. More preferably it is 0.002% or more.
  • Ni is made 0.10% or less. Preferably, it is 0.09% or less.
  • Mg is an element which can control the form of sulfides by addition in a trace amount. If less than 0.0001%, the effect of addition is not sufficiently obtained, so Mg preferably is 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 is an element which can control the form of sulfides by addition in a trace amount. If less than 0.001%, the effect of addition is not sufficiently obtained, so Ca preferably is made 0.001% or more. More preferably it is 0.003% or more.
  • Ca is made 0.050% or less.
  • Y like Mg and Ca, is an element which can control the form of sulfides by addition in a trace amount. If less than 0.001%, the effect of addition is not sufficiently obtained, so Y preferably is made 0.001% or more. More preferably it is 0.003% or more.
  • Y is made 0.050% or less. Preferably it is 0.031% or less.
  • Zr like Mg, Ca, and Y, is an element which can control the form of sulfides by addition in a trace amount. If less than 0.001%, the effect of addition is not sufficiently obtained, so Zr preferably is made 0.001% or more. More preferably it is 0.004% or more.
  • Zr is made 0.050% or less. Preferably it is 0.045% or less.
  • La is an element effective for control of the form of sulfides by addition in a trace amount. Further, it is an element which segregates at the grain boundaries and lowers the number ratio of carbides at the grain boundaries. If less than 0.001%, the effect of control of the form is not sufficiently obtained, so La is preferably made 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 able to control the form of sulfides by addition in a trace amount. Further, it is an element which segregates at the grain boundaries and lowers the number ratio of carbides at the grain boundaries. If less than 0.001%, the effect of control of the form is not sufficiently obtained, so Ce is preferably made 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 remainder of the chemical composition of the steel plate of the present invention is comprised of Fe and unavoidable impurities.
  • the structure of the steel plate of the present invention is substantially a structure comprised of ferrites and carbides.
  • the carbides include cementite (Fe 3 C) which is a compound of iron and carbon, a compound obtained by substituting Mn, Cr, etc. for the Fe atoms in the cementite, and alloy carbides (M 23 C 6 , M 6 C, MC, etc., where M is Fe and other metal elements).
  • a shear zone is formed at the macrostructure of the steel plate and slip deformation occurs concentrated near the shear zone.
  • slip deformation along with proliferation of dislocations, a region of a high dislocation density is formed near the shear zone.
  • slip deformation is promoted and the dislocation density increases.
  • formation of a shear zone can be understood as the phenomenon of slip occurring at a certain one grain crossing the crystal grain boundary and being continuously propagated to the adjoining grain. Accordingly, to suppress the formation of a shear zone, it is necessary to prevent propagation of slip crossing crystal grain boundaries.
  • the carbides in steel plate are strong particles inhibiting slip. By forming carbides at the ferrite grain boundaries, it becomes possible to suppress the formation of a shear zone and improve the cold forgeability.
  • the average particle size of carbides is made 0.4 ⁇ m to 2.0 ⁇ m. If the particle size of the carbides is less than 0.4 ⁇ m, the steel plate remarkably increases in hardness and the cold forgeability falls. More preferably it is 0.6 ⁇ m or more.
  • the average particle size of the carbides exceeds 2.0 ⁇ m, at the time of cold forming, the carbides form starting points of fractures. More preferably, it is 1.95 ⁇ m or less.
  • cementite a carbide of iron
  • cementite a carbide of iron
  • the steel becomes hard and brittle. Therefore, pearlite has to be reduced as much as possible.
  • the area ratio is made 6% or less.
  • Pearlite has a unique lamellar structure, so can be discerned by observation by an SEM or optical microscope. By calculating the region of the lamellar structure at any cross-section, the area ratio of the pearlite can be found.
  • cold forgeability is considered to be strongly affected by the rate of coverage of the ferrite grain boundaries by carbides.
  • High precision measurement is sought, but measurement of the rate of coverage of ferrite grain boundaries by carbides in a three-dimensional space requires serial sectioning SEM observation using an FIB to repeatedly cut a sample for observation in a scanning electron microscope or 3D EBSP observation. A massive measurement time is required and technical knowhow has to be built up.
  • the inventors clarified this and searched for a simpler, higher precision indicator for evaluation and as a result discovered that it is possible to evaluate the cold forgeability by using the ratio of the number of carbides at the ferrite grain boundaries to the number of carbides in the ferrite grains as an indicator and that if the ratio of the number of carbides at the ferrite grain boundaries to the number of carbides in the ferrite grains exceeds 1, the cold forgeability remarkably rises.
  • the carbides are observed by a scanning electron microscope. Before observation, the sample for observation of the structure is polished by wet polishing by Emery paper and a diamond abrasive having an average particle size of 1 ⁇ m, the observed surface is polished to a mirror finish, then a saturated picric acid-alcohol solution is used to etch the structure.
  • the magnification of the observation was made 3000X and images of eight fields of 30 ⁇ m ⁇ 40 ⁇ m at a plate thickness 1/4 layer were captured at random.
  • the obtained structural images were analyzed by image analyzing software such as one made by Mitani Shoji (Win ROOF) to measure in detail the areas of the carbides contained in those regions.
  • the number of carbides which present at the ferrite grain boundaries are counted, the number of carbides at the grain boundaries are subtracted from the total number of carbides, and the number of carbides in the ferrite grains are found. Based on the measured number, the number ratio of carbides at the grain boundaries with respect to the carbides inside the ferrite grains is calculated.
  • the ferrite grain 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 is measured by using the above-mentioned procedure to polish the observed surface of the sample for observation of structure to a mirror finish, then observing the structure of the observed surface etched by a 3% nitric acid-alcohol solution by an optical microscope or scanning electron microscope and applying the line segment method to the captured image.
  • the Vickers hardness of the steel plate 100HV to 180HV it is possible to improve the cold forgeability and impact resistance characteristic after carburizing, quenching, and tempering. If the Vickers hardness is less than 100HV, buckling easily occurs during cold forging, folding and twisting occur at the buckled part, and the impact resistance characteristic falls, so the Vickers hardness is made 100HV or more. Preferably it is 110HV or more.
  • the Vickers hardness exceeds 180HV, the ductility falls, internal cracking easily occurs at the time of cold forging, and the impact resistance characteristic deteriorates, so the Vickers hardness is made 180HV or less. Preferably it is 170HV or less.
  • FIGS. 1A to 1C schematically show an outline of the cold forging test and form of a crack introduced by cold forging.
  • FIG. 1A shows a disk-shaped test material cut out from a hot rolled steel plate
  • FIG. 1B shows the shape of a test material after cold forging
  • FIG. 1C shows the cross-sectional shape of a cold forged test material.
  • FIGS. 1A to 1C from a plate thickness 5.2 mm hot rolled steel plate, a diameter 70 mm disk-shaped test material 1 was cut out (see FIG. 1A ) and deep drawn so as to prepare a cup-shaped test material with a bottom surface of a diameter of 30 mm (not shown).
  • the thickened cup-shaped test material 2 is cut by a wire cut electrical discharge machine made by FANUC so that the cross-section of the diameter part appeared (see FIG. 1C .
  • the cut surface is polished to a mirror finish and the presence of a fracture 3 at the cut surface was confirmed.
  • the hot rolled steel plate covered by the present invention is not limited to a plate thickness 5.2 mm hot rolled steel plate.
  • the present invention can improve the cold forgeability and the impact resistance characteristic after carburizing, quenching, and tempering even in a general plate thickness (2 to 15 mm) hot rolled steel plate.
  • the technical idea of the method of production of the present invention is to integrally manage the rolling conditions and annealing conditions when producing steel plate from a steel slab of the above-mentioned chemical composition so as to improve the cold forgeability and the impact resistance characteristic after carburizing, quenching, and tempering.
  • Molten steel having the required chemical composition is continuously cast into a slab.
  • the slab is used for hot rolling as is in accordance with an ordinary method or is cooled once, then heated and used for hot rolling.
  • the finish hot rolling is ended in the 650°C to 950°C temperature region.
  • the hot rolled steel plate after finish rolling is cooled on the ROT and coiled by a coiling temperature of 400°C to 600°C.
  • the hot rolled steel plate is pickled, then held at two temperature regions as two-stage step type annealing, but at that time, in the first stage annealing, the hot rolled steel plate is heated until the annealing temperature by a 30°C/hour to 150°C/hour heating rate and held at a 650°C to 720°C temperature region for 3 hours to 60 hours for annealing.
  • the hot rolled steel plate is heated until the annealing temperature by a 1°C/hour to 80°C/hour heating rate and held at a 725°C to 790°C temperature region for 3 hours to 50 hours for annealing.
  • the annealed hot rolled steel plate is cooled down to 650°C by a cooling rate of 1°C/hour to 100°C/hour, then is cooled down to room temperature.
  • Molten steel having the required chemical composition is continuously cast into a slab.
  • the slab is used for hot rolling as is or cooled once, then heated.
  • the finish hot rolling is ended in the 650°C to 950°C temperature region.
  • the hot rolled steel plate is coiled at 400°C to 600°C.
  • the slab heating temperature is preferably 1300°C or less, while the heating time where the slab is held at a temperature of the slab surface layer of 1000°C or more is preferably 7 hours or less.
  • the heating temperature exceeds 1300°C or the heating time exceeds 7 hours, the decarburization of the slab surface layer becomes remarkable.
  • 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, while the heating time is 6 hours or less.
  • the finish hot rolling is ended at 650°C to 950°C in temperature. If the finish hot rolling temperature is less than 650°C, the rolling load remarkably rises due to the increase of the deformation resistance of the steel material. Furthermore, the amount of roll wear increases and the productivity falls, so the finish hot rolling temperature is made 650°C or more. Preferably it is 680°C or more.
  • the finish hot rolling temperature exceeds 950°C, bulky scale is formed during passage through the ROT (Run Out Table). Due to the scale, flaws form at the surface of the steel plate.
  • the finish hot rolling temperature is made 950°C or less. Preferably it is 920°C or less.
  • the cooling rate when cooling the hot rolled steel plate on the ROT is preferably 10°C/sec to 100°C/sec. If the cooling rate is less than 10°C/sec, in the middle of the cooling, it is not possible to suppress the formation of bulky scale and the formation of flaws due to the same and the impact resistance characteristic falls, so the cooling rate is preferably 10°C/sec or more. More preferably it is 20°C/sec or more.
  • the outermost layer part is excessively cooled and bainite, martensite, and other low temperature transformed structures are formed at the outermost layer part.
  • the cooling rate is preferably 100°C/sec or less. More preferably it is 80°C/sec or less.
  • the cooling rate indicates the cooling ability received from the cooling facilities in a water spray section at the time when being cooled on the ROT down to the target temperature of coiling from the time when the hot rolled steel plate after finish hot rolling is water cooled at a water spray section after passing through a non-water spray section. It does not show the average cooling rate from the starting point of water spray to the temperature at which the steel plate is coiled up by the coiler.
  • the coiling temperature is made 400°C to 600°C. This is a temperature lower than the general coiling temperature.
  • the structure of the steel plate can be made a bainite structure in which carbides are dispersed in fine ferrite.
  • the coiling temperature is made 400°C or more. Preferably, it is 430°C or more.
  • the coiling temperature exceeds 600°C, pearlite with a large lamellar spacing is formed, high thermal stability bulky needle shaped carbides are formed, and needle shaped carbides remain even after two-stage step type annealing. At the time of cold forging, cracks occur and grow starting from these needle shaped carbides, so the coiling temperature is made 600°C or less. Preferably it is 570°C or less.
  • the hot rolled steel plate produced under the above conditions was pickled, then held in two temperature regions for two-stage step type annealing.
  • the carbides are controlled in stability and the formation of carbides at the ferrite grain boundaries is promoted.
  • the carbides are made to coarsen and added metal elements are made to concentrate to raise the thermal stability of the carbides. After that, the steel is raised to the Ac1 or more in temperature to form austenite in the structure, the fine carbides in the ferrite grains are made to dissolve into the austenite, and coarse carbides are left in the austenite.
  • the austenite is transformed to ferrite and raises the concentration of carbon in the austenite.
  • carbon atoms are adsorbed at the carbides remaining in the austenite, the carbides and austenite cover the grain boundaries of the ferrite, and, finally, it becomes possible to form a structure with a large amount of carbides present at the ferrite grain boundaries. For this reason, it is clear that the structure of the present invention cannot be formed by just simple annealing.
  • the heating rate up to the first stage annealing temperature is made 30°C/hour to 150°C/hour. If the heating rate is less than 30°C/hour, time is required for raising the temperature and the productivity falls, so the heating rate is made 30°C/hour or more. Preferably, it is 40°C/hour or more.
  • the heating rate is over 150°C/hour, the temperature difference between the outer circumferential part and the inside part of the coil increases, scratches and seizing occur due to the difference in heat expansion, and relief shapes are formed at the steel plate surface.
  • the heating rate is made 150°C/hour or less. Preferably, it is made 120°C/hour or less.
  • the annealing temperature in the first stage annealing (first stage annealing temperature) is made 650°C to 720°C. If the first stage annealing temperature is less than 650°C, the carbides becomes insufficient in stability and it becomes difficult to form carbides remaining in the austenite in the second stage annealing, so the first stage annealing temperature is made 650°C or more. Preferably it is 670°C or more.
  • the annealing temperature is made 720°C or less. Preferably it is 700°C or less.
  • the holding time in the first stage annealing (first stage holding time) is made 3 hours to 60 hours. If the first stage holding time is less than 3 hours, the carbides become insufficient in stability and it becomes difficult to form carbides remaining in the second stage annealing, so the first stage holding time is made 3 hours or more. Preferably it is 10 hours or more.
  • the first stage holding time exceeds 60 hours, no further improvement in stability of the carbides can be expected. Furthermore, a drop in productivity is invited, so the first stage holding time is made 60 hours or less. Preferably it is 50 hours or less.
  • the hot rolled steel plate After finishing being held at the first stage annealing, the hot rolled steel plate is heated up to the annealing temperature by a heating rate of 1°C/hour to 80°C/hour. If cooling without performing this second stage annealing, the ferrite grain size does not become larger and the ideal structure cannot be obtained.
  • the heating rate at the second stage annealing is preferably small.
  • the heating rate is less than 1°C/hour, time is required for raising the temperature and the productivity falls, so the heating rate is made 1°C/hour or more. Preferably it is 10°C/hour or more.
  • the heating rate exceeds 80°C/hour, the temperature difference between the outer circumferential part and inside part of the coil increases. Due to the large difference in heat expansion due to deformation, scratches and seizing occur and relief shapes are formed at the surface of the steel plate. At the time of cold forging, cracks form starting from the relief shapes leading to a drop in cold forgeability and a drop in impact resistance characteristic after carburizing, quenching, and tempering, so the heating rate is made 80°C/hour or less.
  • the annealing temperature in the second stage annealing (second stage annealing temperature) is made 725°C to 790°C. If the second stage annealing temperature is less than 725°C, the amount of production of austenite becomes smaller. After cooling after the second stage annealing, the number ratio of carbides at the ferrite grain boundaries falls and, further, the ferrite grain size becomes smaller. For this reason, the second stage annealing temperature is made 725°C or more. Preferably it is 735°C or more.
  • the second stage annealing temperature exceeds 790°C, it becomes difficult to form carbides remaining in the austenite and control to the above-mentioned change of structure becomes difficult, so the second stage annealing temperature is made 790°C or less. Preferably it is 780°C or less.
  • the holding time in the second stage annealing is made 1 hour to 50 hours. If the second stage holding time is less than 1 hour, the amount of austenite which is produced is small, the carbides in the ferrite grains are not sufficiently dissolved, it becomes difficult to increase the number ratio of carbides at the ferrite grain boundaries, and, further, the ferrite grains become smaller in size, so the second stage holding time is made 1 hour or more. Preferably, it is 5 hours or more.
  • the second stage holding time exceeds 50 hours, it is difficult to make carbides remain in the austenite, so the second stage holding time is made 50 hours or less. Preferably it is 45 hours.
  • the annealed hot rolled steel plate is gradually cooled down to 650°C by a cooling rate of 1°C/hour to 100°C/hour. If the stop temperature of gradual cooling exceeds 650°C, due to the cooling rate subsequently exceeding 100°C/hour down to room temperature, nontransformed austenite transforms to pearlite or bainite, the hardness increases, and the cold forgeability falls, so the cooling stop temperature is made 650°C.
  • a slower cooling rate is preferable. If the cooling rate is less than 1°C/hour, the time required for cooling increases and the productivity falls, so the cooling rate is made 1°C/hour or more. Preferably it is 10°C/hour or more.
  • the cooling rate exceeds 100°C/hour, austenite transforms to pearlite and the steel plate increases in hardness so a drop in cold forgeability and a drop in impact resistance characteristics after carburizing, quenching, and tempering are invited, so the cooling rate is made 100°C/hour or less. Preferably it is 90°C/hour.
  • the cooling stop temperature is the temperature where the cooling rate should be used for control. If cooling down to 650°C by a cooling rate of 1°C/hour to 100°C/hour, the cooling down to 650°C or less is not particularly limited.
  • the atmosphere of the annealing is not limited to any specific atmosphere.
  • it may be any of an atmosphere of 95% or more of nitrogen, an atmosphere of 95% or more of hydrogen, and an air atmosphere.
  • a continuously cast slab (steel ingot) having a chemical composition shown in Table 1 was heated at 1240°C for 1.8 hours, then was used for hot rolling.
  • the finish hot rolling was ended at 890°C, the steel was cooled on a ROT by a 45°C/sec cooling rate down to 520°C and was coiled up at 510°C to produce a hot rolled coil with a plate thickness of 5.2 mm.
  • the hot rolled coil was pickled, the coil was loaded into a box-type annealing furnace, the atmosphere was controlled to 95% hydrogen-5% nitrogen, then the coil was heated from room temperature up to 705°C by a heating rate of 100°C/hour, was held at 705°C for 36 hours to make the temperature distribution inside the coil uniform, then was heated by a 5°C/hour heating rate up to 760°C, and, furthermore, was held at 760°C for 10 hours, then was cooled down to 650°C by a 10°C/hour cooling rate, then was furnace cooled down to room temperature to prepare a sample for evaluation of the characteristics.
  • the structure of the sample was observed by the above-mentioned method.
  • the crack length at the sample after cold forging was measured by the above-mentioned method.
  • the carburization of the thickened sample was performed by gas carburization. To make the carbon disperse from the inside of the furnace atmosphere gas through the surface layer of the sample to the inside of the steel, the sample was treated by holding it at 940°C for 120 minutes inside a furnace controlled to a carbon potential of 0.5 mass% C, then was furnace cooled down to room temperature.
  • the sample was heated from room temperature to 840°C, then was held for 20 minutes and quenched in 60°C oil.
  • the hardened sample was held at 170°C for 60 minutes, then air cooled for tempering to prepare a carburized, quenched, and tempered sample.
  • FIG. 2 schematically shows an outline of the drop weight test for evaluating the impact resistance characteristic of a carburized, quenched, and tempered sample.
  • the bottom of the cup of a carburized, quenched, and tempered cup-shaped sample 4 was fastened by a fixture.
  • a weight 2 kg dropping weight top side width: 50 mm, bottom side width: 10 mm, height: 80 mm, and length: 110 mm
  • the sample was examined for the presence of any cracking and was evaluated for the impact resistance characteristic.
  • a sample with no fracture or breakage observed as a result of free dropping was evaluated as excellent in impact resistance characteristics, that is, "OK", while a sample with a fracture or breakage observed was evaluated as inferior in impact resistance, that is, "NG”.
  • Table 2 shows the results of measurement and results of evaluation of the carbide size, pearlite area ratio, ferrite grain size, Vickers hardness, ratio of the number of carbides at the ferrite grain boundaries to number of carbides in the ferrite grains, ratio of maximum crack length to plate thickness at the vertical wall parts, and impact resistance characteristic in the prepared samples.
  • Table 2 Carbide size ( ⁇ m) Ferrite grain size ( ⁇ m) Pearlite area rate (%) Vickers hardness (HV) No. of carbides at grain boundaries/No.
  • Comparative Steel L-1 the amount of C is low and the hardness before cold forging is less than 100HV, so the cold forgeability is low.
  • Comparative Steels M-1, P-1, and Z-1, P, Al, and N are excessively contained and, at the second stage annealing, the amount of segregation at the ⁇ / ⁇ interfaces is large, so formation of carbides at the grain boundaries is suppressed.
  • Comparative Steel S-1 Si is excessively contained and ductility of the ferrite is low, so the cold forgeability is low.
  • Mo and Cr are excessively contained, so carbides finely disperse inside the ferrite grains and the hardness exceeds 180HV.
  • Mn is excessively contained, so the impact resistance characteristic after carburizing, quenching, and tempering is remarkably low.
  • Comparative Steel O-1 the amount of Cr is small and the austenite grains at the surface layer abnormally coarsen at the time of carburization, so the impact resistance is low.
  • Comparative Steel R-1 S is excessively contained, so coarse MnS is formed in the steel and the cold forgeability is low.
  • Comparative Steel U-1 C is excessively contained, so coarse carbides form inside the thickened layer of the steel and coarse carbides remain even after the carburizing and quenching, so the impact resistance characteristic is low.
  • Comparative Steel V-1 the amount of Mn is small and the carbides are hard to raise in stability, so the cold forgeability and impact resistance characteristic after carburizing, quenching, and tempering are low.
  • Comparative Steels W-1 and X-1 O and Ti are excessively contained, so the oxides and TiC present in the ferrite grains become site for formation of carbides in gradual cooling after dual phase region annealing, the formation of carbides at the grain boundaries is suppressed, and the cold forgeability is low.
  • Comparative Steel Y-1 B is excessively contained, so the cold forgeability is low.
  • slabs having the chemical compositions A, B, C, D, E, F, G, H, I, J, and K shown in Table 1 were hot rolled and annealed under the conditions shown in Table 3 to prepare annealed samples of hot rolled plates of a thickness of 5.2 mm.
  • Table 4 shows the results of measurement and results of evaluation of the carbide size, pearlite area ratio, ferrite grain size, Vickers hardness, ratio of the number of carbides at the ferrite grain boundaries to number of carbides in the ferrite grains, ratio of maximum crack length to plate thickness at the vertical wall parts, and impact resistance characteristics in the prepared samples.
  • Table 3 Hot rolling conditions Annealing conditions Remarks Heating temp. (°C) Soaking time (hours) Finish hot rolling temp. (°C) ROT cooling rate (°C/ sec) Coiling temp. (°C) 1st stage 2nd stage Cooling rate down to 650°C (°C/hour) Heating rate (°C/ hour) Holding temp.
  • Comparative Steel E-3 the finish hot rolling temperature is low, the rolling load increases, and the productivity falls.
  • Comparative Steel D-2 the finish hot rolling temperature is high and scale flaws form at the surface of the steel plate, so when subjected to a wear resistance test after quenching and tempering, fractures and peeling occur starting from the scale flaws and the wear resistance characteristic falls.
  • Comparative Steel F-2 the cooling rate at the ROT (Run Out Table) is slow and a drop of productivity and formation of scale flaws are invited.
  • Comparative Steel C-4 the cooling rate at the ROT is 100°C/sec and the outermost layer part of the steel plate is excessively cooled, so fine cracks formed at the outermost layer part.
  • Comparative Steel C-2 the coiling temperature is low, large amounts of bainite, martensite, and other low temperature transformed structures are formed causing embrittlement, fractures frequently form at the time of pay out from the hot rolled coil, and the productivity falls. Furthermore, in a sample taken from a cracked piece, the wear resistance characteristic is low.
  • Comparative Steel G-2 the coiling temperature is high, bulky pearlite of lamellar spacing is formed in the hot rolled structure, the needle shaped coarse carbides become high in thermal stability, and the above carbides remain in the steel plate even after two-stage step type annealing, so the machinability is low.
  • Comparative Steel H-4 the heating rate in the first stage annealing of the two-stage step type annealing is slow, so the productivity is low.
  • the heating rate in the first stage annealing is fast, so the temperature difference between the inside part and inside and outside circumferential parts of the coil becomes larger, scratches and seizing occur due to the difference in thermal expansion, and, when used for evaluating and testing the wear resistance characteristic after quenching and tempering, cracks and peeling occur from the flaw parts and the wear resistance characteristic falls.
  • the holding temperature in the first stage annealing (annealing temperature) is low, the coarsening treatment of carbides at the Ac1 point or less is insufficient, and the carbides are insufficient in thermal stability, so carbides remaining at the second stage annealing are reduced and pearlite transformation in the structure after gradual cooling cannot be suppressed, so the machinability is low.
  • Comparative Steel D-4 the holding temperature in the first stage annealing (annealing temperature) is high, austenite is formed during the annealing, and the carbides cannot be raised in stability, so pearlite is formed after annealing, the Vickers hardness exceeds 180HV, and the machinability is low.
  • the holding time in the first stage annealing is short and the stability of carbides cannot be raised, so the machinability is low.
  • Comparative Steel F-2 the holding time in the first stage annealing is long, the productivity is low, and further seizing flaws occur and the wear resistance characteristic is low.
  • Comparative Steel B-4 the heating rate in the second stage annealing in the two-stage step type annealing is slow, so the productivity is low.
  • Comparative Steel A-3 the heating rate at the second stage annealing is fast, so the temperature difference between the inside part and outer circumferential part of the coil become greater, scratches and seizing occur due to the large difference in heat expansion due to deformation, and the wear resistance characteristic after quenching and tempering is low.
  • the holding temperature in the second stage annealing is low, the amount of production of austenite is small, and the ratio of number of carbides at the ferrite grain boundaries cannot be increased, so the machinability is low.
  • the holding temperature at the second stage annealing is high and dissolution of the carbides during the annealing is promoted, so it becomes difficult to form carbides at the grain boundaries after the gradual cooling and further pearlite is produced, the Vickers hardness exceeds 180HV, and the machinability is low.
  • Comparative Steel J-3 the holding time at the second stage annealing is long and dissolution of the carbides is promoted, so the machinability is low.
  • Comparative Steel D-3 the cooling rate from second stage annealing to 650°C is slow, the productivity is low, coarse carbides are formed in the structure after gradual cooling, cracks are formed starting from the coarse carbides at the time of cold forging, and the cold forgeability falls.
  • the cooling rate from second stage annealing to 650°C is fast, the pearlite transformation occurs at the time of cooling, and the hardness increases, so the cold forgeability is low.
  • the hot rolled coils were pickled, the hot rolled coils were loaded into a box-type annealing furnace, the atmosphere was controlled to 95% hydrogen-5% nitrogen, then the coils were heated from room temperature up to 705°C by a heating rate of 100°C/hour, were held at 705°C for 36 hours to make the temperature distribution inside the coils uniform, then were heated by a 5°C/hour heating rate up to 760°C, and, furthermore, were held at 760°C for 10 hours, then were cooled down to 650°C by a 10°C/hour cooling rate, then were furnace cooled down to room temperature to prepare samples for evaluation of the characteristics.
  • Table 7 shows the results of measurement and results of evaluation of the carbide size, pearlite area ratio, ferrite grain size, Vickers hardness, ratio of the number of carbides at the ferrite grain boundaries to number of carbides in the ferrite grains, ratio of maximum crack length to plate thickness at the vertical wall parts, and impact resistance characteristics in the prepared samples.
  • Table 7 Carbide size ( ⁇ m) Ferrite grain size ( ⁇ m) Pearlite area rate (%) Vickers hardness (HV) No. of carbides at t grain boundaries/ No.
  • Comparative Steels AR-1, AS-1, AW-1, AZ-1, BB-1, and BF-1 La, As, Cu, Ni, Sb, and Ce are excessively contained and the amount of segregation at the ⁇ / ⁇ interface becomes greater at the time of second stage annealing, so formation of carbides at the grain boundaries is suppressed.
  • Comparative Steel BG-1 Si is excessively contained and the ductility of the ferrite is low, so the cold forgeability is low.
  • Comparative Steels AT-1, AV-1, BA-1, BC-1, BH-1, and BJ-1 Mo, Nb, Cr, Ta, W, and V are respectively excessively contained, so carbides finely disperse inside the ferrite grains and the hardness exceeds 180HV.
  • Mo, Nb, Cr, Ta, W, and V are respectively excessively contained, so carbides finely disperse inside the ferrite grains and the hardness exceeds 180HV.
  • Comparative Steel BF-1 Mn is excessively contained, so the impact resistance characteristic after carburizing, quenching, and tempering is remarkably low.
  • Comparative Steels AU-1, AX-1, AY-1, and BE-1 Zr, Ca, Mg, and Y are respectively excessively contained, coarse oxides or nonmetallic inclusions are formed in the steel, cracks form starting from the coarse oxides or coarse nonmetallic inclusions at the time of cold forging, and the cold forgeability falls.
  • Comparative Steel BD-1 Sn is excessively contained, the ferrite becomes brittle, and the cold forgeability is low.
  • Comparative Steel BK-1 C is excessively contained, so coarse carbides form at the inside of the increased thickness part of the steel, coarse carbides remain even after carburizing and quenching, and the impact resistance characteristic also falls.
  • slabs having the chemical compositions of AA, AB, AC, AD, AE, AF, AG, AH, AI, AJ, AK, AL, AM, AN, AO, AP, and AQ shown in Table 5 were hot rolled and annealed under the conditions shown in Table 8 to fabricate annealed samples of hot rolled plates of thicknesses of 5.2 mm.
  • Table 8 Hot rolling conditions Annealing conditions Remarks Heating temp. (°C) Soaking time (hours) Finish hot rolling temp. (°C) ROT cooling rate (°C/ sec) Coiling temp.
  • Table 9 shows the results of measurement and results of evaluation of the carbide size, pearlite area ratio, ferrite grain size, Vickers hardness, ratio of the number of carbides at the ferrite grain boundaries to number of carbides in the ferrite grains, ratio of maximum crack length to plate thickness at the vertical wall parts, and impact resistance characteristic in the prepared samples.
  • Table 9 Carbide size ( ⁇ m) Ferrite grain size ( ⁇ m) Pearlite area rate (%) Vickers hardness (HV) No. of carbides at grain boundaries/ No.
  • Comparative Steel AC-2 the finish hot rolling temperature is low and the productivity is low.
  • Comparative Steel AN-4 the finish hot rolling temperature is high, scale flaws form at the surface of the steel plate and cracks form from the flaw parts when impact load was given after cold forging and carburizing, quenching, and tempering, and the impact resistance characteristic falls.
  • Comparative Steel AN-3 the coiling temperature is low, large amounts of bainite, martensite, and other low temperature transformed structures are produced resulting in embrittlement, fractures frequently occur at the time of pay out of the hot rolled coil, and the productivity falls. Furthermore, at the sample taken from the cracked slab, the cold forging and impact resistance characteristic after carburizing, quenching, and tempering are inferior.
  • Comparative Steel AH-3 the coiling temperature is high, bulky pearlite of the lamellar spacing is formed in the hot rolled structure, needle-shaped coarse carbides are high in thermal stability, and even after two-stage step type annealing, the above carbides remain in the steel plate, so the cold forgeability is low.
  • Comparative Steel AF-4 the heating rate in the first stage annealing of the two-stage step type annealing is slow, so the productivity is low.
  • Comparative Steel AG-2 the heating rate in the first stage annealing is fast, so the difference in temperature between the inside part and outer circumferential part of the coil becomes larger, scratches and seizing due to the difference in heat expansion occur, and the cold forging and impact resistance characteristic after carburizing, quenching, and tempering fall.
  • the holding temperature in the first stage annealing (annealing temperature) is low, the coarsening of the carbides at the Ac1 point or less is insufficient, the thermal stability of the carbides becomes insufficient, the carbides remaining at the time of the second stage annealing decrease, the pearlite transformation cannot be suppressed in the structure after gradual cooling, and the cold forgeability falls.
  • Comparative Steel AM-3 the first stage holding temperature (annealing temperature) is high, austenite is produced during the annealing, the stability of the carbides cannot be raised, and the cold forgeability and impact resistance characteristic after carburizing, quenching, and tempering fall.
  • Comparative Steel AF-2 the holding time in the first stage annealing is short, the stability of the carbides cannot be raised, and the cold forgeability is low.
  • Comparative Steel AO-4 the holding time in the first stage annealing is long and the productivity is low.
  • Comparative Steel AP-4 the heating rate at the second stage annealing in the two-stage step type annealing is slow, so the productivity is low.
  • Comparative Steel AI-3 the heating rate at the second stage annealing is fast, so the temperature difference between the inside part and the outer circumferential part of the coil become greater, scratches and seizing occur due to the large difference in heat expansion due to transformation, and, when an impact load is given after carburizing, quenching, and tempering, fractures occur from the flaw parts and the impact resistance characteristics fall.
  • the holding temperature in the second stage annealing is low, the amount of production of austenite is small, it is not possible to increase the number ratio of carbides at the ferrite grain boundaries, and the cold forgeability falls.
  • the holding temperature in the second stage annealing is high, the dissolution of carbides during annealing is promoted, and therefore it becomes difficult to cause the production of carbides at the grain boundaries after gradual cooling, and the cold forgeability and impact resistance characteristics after carburizing, quenching, and tempering fall.
  • Comparative Steel AJ-4 the holding time in the second stage annealing is long and dissolution of carbides is promoted, so the cold forgeability is low.
  • Comparative Steel AQ-3 the cooling rate from the second stage annealing to 650°C is slow so the productivity is low and coarse carbides are formed in the structure after gradual cooling so cracks formed starting from the coarse carbides at the time of cold forging and the cold forgeability dropped.
  • Comparative Steel AP-2 the cooling rate from the second stage annealing to 650°C was slow, pearlite transformation occurred at the time of cooling, and the hardness increased, so the cold forgeability fell.
  • FIG. 3 shows the relationship among the ratio of the number of carbides at the grain boundaries to the number of carbides in the grains, and the crack length and impact resistance characteristics of cold forging test pieces after carburizing, quenching, and tempering.
  • FIG. 4 shows another relationship between the ratio of the number of carbides at the grain boundaries to the number of carbides in the grains and the crack length of cold forging test pieces and impact resistance characteristic after carburizing, quenching, and tempering.
  • FIG. 4 is a view showing that it is possible to keep down crack length even in steel plate to which additional elements are added.
  • the steel plate of the present invention is, for example, suitable as a material when forming a part by cold forging such as plate working to obtain a high cycle gear or other part, so the present invention has high industrial applicability.
EP16800074.3A 2015-05-26 2016-05-25 Stahlblech und verfahren zur herstellung davon Withdrawn EP3305929A4 (de)

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