WO2020158357A1 - High-carbon hot-rolled steel sheet and method for manufacturing same - Google Patents

High-carbon hot-rolled steel sheet and method for manufacturing same Download PDF

Info

Publication number
WO2020158357A1
WO2020158357A1 PCT/JP2020/000783 JP2020000783W WO2020158357A1 WO 2020158357 A1 WO2020158357 A1 WO 2020158357A1 JP 2020000783 W JP2020000783 W JP 2020000783W WO 2020158357 A1 WO2020158357 A1 WO 2020158357A1
Authority
WO
WIPO (PCT)
Prior art keywords
less
steel sheet
rolled steel
cementite
hot
Prior art date
Application number
PCT/JP2020/000783
Other languages
French (fr)
Japanese (ja)
Inventor
友佳 宮本
櫻井　康広
義彦 小野
Original Assignee
Ｊｆｅスチール株式会社
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Ｊｆｅスチール株式会社 filed Critical Ｊｆｅスチール株式会社
Priority to US17/425,824 priority Critical patent/US20220170126A1/en
Priority to KR1020217023358A priority patent/KR102570145B1/en
Priority to EP20749360.2A priority patent/EP3901303A4/en
Priority to CN202080011346.5A priority patent/CN113366137B/en
Priority to JP2020520327A priority patent/JP6927427B2/en
Publication of WO2020158357A1 publication Critical patent/WO2020158357A1/en

Links

Classifications

    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/004—Heat treatment of ferrous alloys containing Cr and Ni
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties 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
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0273—Final recrystallisation annealing
    • C—CHEMISTRY; METALLURGY
    • 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
    • C22C38/008—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
    • C22C38/20—Ferrous alloys, e.g. steel alloys containing chromium 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
    • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—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
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel 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
    • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/003—Cementite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/009—Pearlite

Definitions

  • the present invention relates to a high carbon hot-rolled steel sheet excellent in cold workability and hardenability (dip hardenability and carburizing hardenability) and a method for producing the same.
  • high carbon hot-rolled steel sheets are carbon steel steels for machine structures and alloy steels for machine structures specified in JIS G4051. After being processed into a desired shape, it is often manufactured by quenching to secure a desired hardness. Therefore, the hot-rolled steel sheet used as a material is required to have excellent cold workability and hardenability, and various steel sheets have been proposed so far.
  • Patent Document 1 C: 0.15 to 0.9%, Si: 0.4% or less, Mn: 0.3 to 1.0%, P: 0.03% or less in weight%.
  • T. Al: 0.10% or less, Cr: 1.2% or less, Mo: 0.3% or less, Cu: 0.3% or less, Ni: 2.0% or less, or Ti: 0. 01 to 0.05%, B: 0.0005 to 0.005%, N: 0.01% or less, a spheroidization rate of 80% or more, and an average particle size of 0.4 to 1.0 ⁇ m.
  • a high carbon steel sheet for precision punching is described which has a structure in which carbides are dispersed in ferrite.
  • the composition is such that C: 0.2% or more, Ti: 0.01 to 0.05%, and B: 0.0003 to 0.005% by mass%, and the average particle diameter of carbide is Describes a high-carbon steel sheet having improved workability in which the ratio of carbides having a grain size of 1.0 ⁇ m or less and 0.3 ⁇ m or less is 20% or less.
  • Patent Document 3 C: 0.20% or more and 0.45% or less, Si: 0.05% or more and 0.8% or less, Mn: 0.5% or more and 2.0% or less, P in mass% : 0.001% to 0.04%, S: 0.0001% to 0.006%, Al: 0.005% to 0.1%, Ti: 0.005% to 0.2% , B: 0.001% or more and 0.01% or less, and N: 0.0001% or more and 0.01% or less, Cr: 0.05% or more and 0.35% or less, Ni: 0.01% or more 1 0.0% or less, Cu: 0.05% or more and 0.5% or less, Mo: 0.01% or more and 1.0% or less, Nb: 0.01% or more and 0.5% or less, V: 0.01% Or more and 0.5% or less, Ta: 0.01% or more and 0.5% or less, W: 0.01% or more and 0.5% or less, Sn: 0.003% or more and 0.03% or less, Sb: 0. A B-added steel having one or more components of 003%
  • Patent Document 4 C: 0.10 to 1.2%, Si: 0.01 to 2.5%, Mn: 0.1 to 1.5%, and P: 0.04% or less in mass%. , S: 0.0005 to 0.05%, Al: 0.2% or less, Te: 0.0005 to 0.05%, N: 0.0005 to 0.03%, and Sb: 0.001 to 0 0.05%, Cr: 0.2-2.0%, Mo: 0.1-1.0%, Ni: 0.3-1.5%, Cu: 1.0% or less, B:0 A mechanical structure with a composition containing at least one of 0.005% or less, a structure mainly composed of ferrite and pearlite, and having a ferrite grain size of 11 or more and improved cold workability and low decarburization. Steel for use is described.
  • Patent Document 5 C: 0.20 to 0.40%, Si: 0.10% or less, Mn: 0.50% or less, P: 0.03% or less, and S: 0.010 in mass%. % Or less, sol. Al: 0.10% or less, N: 0.005% or less, B: 0.0005 to 0.0050%, and one or more of Sb, Sn, Bi, Ge, Te and Se in total. It contains 0.002 to 0.03%, consists of ferrite and cementite, has a microstructure with a cementite density of 0.10 particles/ ⁇ m 2 or less in ferrite grains, and has a hardness of 75 or less in HRB and a total elongation. Is 38% or more, and a high carbon hot rolled steel sheet having improved hardenability and workability is described.
  • Patent Document 6 in mass%, C: 0.20 to 0.48%, Si: 0.10% or less, Mn: 0.50% or less, P: 0.03% or less, S: 0.010. % Or less, sol. Al: 0.10% or less, N: 0.005% or less, B: 0.0005 to 0.0050%, and one or more of Sb, Sn, Bi, Ge, Te and Se in total. It contains 0.002 to 0.03%, is composed of ferrite and cementite, has a microstructure with a cementite density of 0.10 particles/ ⁇ m 2 or less in the ferrite grains, and has a hardness of HRB of 65 or less, and a total hardness of 65 or less. A high carbon hot rolled steel sheet having an elongation of 40% or more and improved hardenability and workability is described.
  • Patent Document 7 C: 0.20 to 0.40%, Si: 0.10% or less, Mn: 0.50% or less, P: 0.03% or less, and S: 0.010 in mass%. % Or less, sol. Al: 0.10% or less, N: 0.005% or less, B: 0.0005 to 0.0050%, and one or more of Sb, Sn, Bi, Ge, Te and Se in total.
  • the content of 0.002 to 0.03%, the proportion of solid solution B in the B content is 70% or more, and it is composed of ferrite and cementite, and the cementite density in the ferrite grains is 0.08 pieces/ ⁇ m 2 or less.
  • a high carbon hot rolled steel sheet having a HRB of 73 or less and a total elongation of 39% or more.
  • Patent Document 8 C: 0.15 to 0.37%, Si: 1% or less, Mn: 2.5% or less, P: 0.1% or less, S: 0.03% or less in mass%.
  • Patent Document 2 pays attention not only to the average grain size of carbides, but also to the fact that fine carbides of 0.3 ⁇ m or less affect workability, and controls the average grain size of carbides to 1.0 ⁇ m or less, In addition, the proportion of carbides of 0.3 ⁇ m or less is controlled to 20% or less. This describes that a steel sheet with improved workability can be obtained, and further, a steel sheet having Ti and B added and having excellent hardenability is described. However, Patent Document 2 does not describe solid solution B or the like that affects the hardenability, and does not describe at which position of the steel plate the hardened hardness corresponds.
  • Patent Document 3 describes that a steel having improved cold workability and decarburization resistance can be obtained by adjusting the composition of components.
  • Patent Document 3 there is no description regarding the dip quenching property and the carburizing quenching property.
  • Patent Document 4 contains B and further one or more components of Cr, Ni, Cu, Mo, Nb, V, Ta, W, Sn, Sb, As, and a solid layer in the surface layer. It is stated that by ensuring a predetermined amount of molten B, a steel that achieves high hardenability can be obtained.
  • the hydrogen concentration in the atmosphere in the annealing step is specified to be 95% or more, and there is no description on whether it is possible to suppress nitrification and secure the solid solution B in the annealing step in the nitrogen atmosphere. ..
  • Patent Documents 5 to 7 have the effect of preventing nitriding by containing 0.002 to 0.03% of B and at least one of Sb, Sn, Bi, Ge, Te and Se in total. It is described that, even when annealing is performed in a nitrogen atmosphere, for example, nitrification is prevented and the solid solution B is maintained at a predetermined amount to enhance the hardenability. However, in any of Patent Documents 5 to 7, there is no description about quenching hardness in the surface layer.
  • Patent Document 8 proposes a steel with high hardenability by containing C: 0.15 to 0.37% and at least one of B, Sb, and Sn. However, Patent Document 8 does not consider higher quenchability such as carburizing quenchability.
  • the present invention has been made in view of the above problems, and provides a high-carbon hot-rolled steel sheet having excellent cold workability and excellent hardenability (dub hardenability, carburizing hardenability) and a method for producing the same. With the goal.
  • cementite having an equivalent circle diameter of 0.1 ⁇ m or less is It has a great influence.
  • tensile strength of 480 MPa or less and total elongation (El) of 33% or more can be obtained.
  • Cementite with a circle equivalent diameter of 0.1 ⁇ m or less greatly affects the hardness (hardness) and total elongation of the high carbon hot rolled steel sheet before quenching.
  • tensile strength of 440 MPa or less and total elongation (El) of 36% or more can be obtained.
  • finish rolling is performed at a finish rolling finish temperature: Ar 3 transformation point or higher, and then cooled to 650 to 750° C. at an average cooling rate of 20 to 100° C./sec, and a winding temperature: 500 to After winding at 700° C., cooling to room temperature to form a hot-rolled steel sheet, the hot-rolled steel sheet is heated at an average heating rate of 15° C./h or more between 450 to 600° C., and an annealing temperature: Ac 1 transformation point.
  • Predetermined microstructure can be secured by annealing for 1.0 h or more.
  • finish rolling is performed: finish rolling at a finish temperature of Ar 3 transformation point or higher, and thereafter, cooling is performed at an average cooling rate of 20 to 100° C./sec to 650 to 750° C., and a winding temperature is: After winding at 500 to 700°C and cooling to room temperature to form a hot rolled steel sheet, the hot rolled steel sheet is heated between 450 and 600°C at an average heating rate of 15°C/h or more to obtain an Ac 1 transformation point or more.
  • Group A Ti: 0.06% or less
  • Group B One or two or more selected from Nb, Mo, Ta, Ni, Cu, V, and W, respectively, 0.0005% or more 0.1 % Or less
  • a high-carbon hot-rolled steel sheet which is annealed by heating the hot-rolled steel sheet at an average heating rate of 15° C./h or more in a temperature range of 450 to 600° C. and holding it at an annealing temperature of less than Ac 1 transformation point for 1.0 hour or more.
  • Manufacturing method [6] The method for producing a high-carbon hot-rolled steel sheet according to any one of [1] to [4], wherein the steel having the above-mentioned composition of ingredients is subjected to hot rough rolling, and then finish rolling finish temperature: Ar 3 transformation.
  • Finish rolling is performed at a point or higher, then the average cooling rate: 20 to 100°C/sec is cooled to 650 to 750°C, and the coiling temperature is 500 to 700°C.
  • the rolled steel sheet is heated to a temperature range of 450 to 600° C. at an average heating rate of 15° C./h or more, and is kept for 0.5 h or more at an Ac 1 transformation point or more and an Ac 3 transformation point or less, and then an average cooling rate: 1 to cooled to below Ar 1 transformation point at 20 ° C. / h, the method of producing a high-carbon hot-rolled steel sheet subjected to annealing for holding 20h or less than Ar 1 transformation point.
  • Component composition The component composition of the high carbon hot-rolled steel sheet of the present invention and the reason for limitation thereof will be described.
  • “%” which is a unit of the content of the following component composition shall mean “mass %” unless there is particular notice.
  • Si 0.8% or less Si is an element that increases strength by solid solution strengthening.
  • the amount of Si is 0.8% or less because it hardens as the amount of Si increases and the cold workability deteriorates. It is preferably 0.65% or less, more preferably 0.50% or less. When further cold workability is required in the use of difficult-to-form parts, it is preferably 0.30% or less. From the viewpoint of ensuring a predetermined softening resistance in the tempering process after quenching, the Si amount is preferably 0.1% or more, more preferably 0.2% or more.
  • Mn 0.10% or more and 0.80% or less Mn is an element that improves hardenability and increases strength by solid solution strengthening. If the Mn content is less than 0.10%, both the quench hardenability and the carburizing hardenability begin to deteriorate, so the Mn content is set to 0.10% or more. In the case of reliably quenching the inside of a thick material or the like, the content is preferably 0.25% or more, more preferably 0.30% or more. On the other hand, when the Mn content exceeds 0.80%, a band structure due to Mn segregation develops, the structure becomes nonuniform, and solid solution strengthens the steel to deteriorate the cold workability. Therefore, the amount of Mn is 0.80% or less. As a material for parts required to have moldability, a predetermined cold workability is required, so that the Mn content is preferably 0.65% or less. More preferably, it is 0.55% or less.
  • P 0.03% or less
  • P is an element that increases strength by solid solution strengthening. If the P content exceeds 0.03%, grain boundary embrittlement is caused, and the toughness after quenching deteriorates. Further, cold workability is also reduced. Therefore, the P content is 0.03% or less. In order to obtain excellent toughness after quenching, the P content is preferably 0.02% or less. Since P reduces the cold workability and the toughness after quenching, the smaller the amount of P, the more preferable. However, if the P content is excessively reduced, the refining cost increases, so the P content is preferably 0.005% or more. More preferably, it is 0.007% or more.
  • S 0.010% or less
  • S is an element that must be reduced because it forms a sulfide and reduces the cold workability and the toughness of the high carbon hot-rolled steel sheet after quenching. If the S content exceeds 0.010%, the cold workability and the toughness of the high carbon hot-rolled steel sheet after quenching are significantly deteriorated. Therefore, the S amount is 0.010% or less.
  • the S content is preferably 0.005% or less. Since S lowers the cold workability and the toughness after quenching, it is preferable that the amount of S is smaller. However, if the S content is excessively reduced, the refining cost increases, so the S content is preferably 0.0005% or more.
  • the N content is 0.01% or less. It is preferably 0.0065% or less. More preferably, it is 0.0050% or less.
  • N forms AlN, a Cr-based nitride, and a B-nitride. This is an element that appropriately suppresses the growth of austenite grains during heating during the quenching treatment and improves the toughness after quenching. Therefore, the N content is preferably 0.0005% or more. More preferably, it is 0.0010% or more.
  • B 0.0005% or more and 0.005% or less
  • B is an important element that enhances hardenability.
  • the amount of B is less than 0.0005%, no sufficient effect is observed, so the amount of B must be 0.0005% or more. It is preferably 0.0010% or more.
  • the B content is more than 0.005%, recrystallization of austenite after finish rolling is delayed, resulting in the development of texture of the hot rolled steel sheet, the anisotropy after annealing becomes large, and in draw forming. Ears are more likely to occur. Therefore, the B content is 0.005% or less. It is preferably 0.004% or less.
  • Total of one or two selected from Sn and Sb are elements effective for suppressing nitriding from the steel sheet surface layer. If the total of one or more of these elements is less than 0.002%, a sufficient effect is not observed, so the total of one or more of these elements is set to 0.002% or more. More preferably, it is 0.005% or more. On the other hand, even if the total content of one or more of these elements exceeds 0.1%, the effect of preventing nitrification is saturated. Further, since these elements tend to segregate at the grain boundaries, if the total content exceeds 0.1%, the content becomes too high, which may cause grain boundary embrittlement. Therefore, the total content of one or two selected from Sb and Sn is 0.1% or less. It is preferably 0.03% or less, more preferably 0.02% or less.
  • the present invention by controlling the total of one or two selected from Sb and Sn to be 0.002% or more and 0.1% or less, it is possible to suppress nitriding from the surface layer of the steel sheet even when annealed in a nitrogen atmosphere.
  • the increase in nitrogen concentration in the surface layer of the steel sheet is suppressed.
  • the present invention since the nitriding from the steel sheet surface layer can be suppressed, the amount of solid solution B in the region from the steel sheet surface layer after annealing to a depth of 100 ⁇ m can be suppressed even when annealed in a nitrogen atmosphere.
  • the balance other than the above is Fe and inevitable impurities.
  • the high-carbon hot-rolled steel sheet of the present invention can obtain the desired characteristics.
  • the high-carbon hot-rolled steel sheet of the present invention may contain the following elements, if necessary, for the purpose of further improving hardenability.
  • Nb, Mo, Ta, Ni, Cu, V, and W may be added in the required amounts, respectively. Good.
  • Mo 0.0005% or more and 0.1% or less Mo is an element effective for improving hardenability and temper softening resistance. If less than 0.0005%, the effect of addition is small. Therefore, when Mo is contained, the lower limit is preferably 0.0005%. More preferably, it is 0.0010% or more. When Mo exceeds 0.1%, the effect of addition is saturated and the cost also increases. Therefore, when Mo is contained, the upper limit is preferably 0.1%. It is more preferably 0.05% or less, and even more preferably less than 0.03%.
  • Ta 0.0005% or more and 0.1% or less Ta forms carbonitrides like Nb, prevents abnormal grain growth of crystal grains during heating before quenching, prevents crystal grain coarsening, and improves temper softening resistance. Is an effective element. If less than 0.0005%, the effect of addition is small. Therefore, when Ta is contained, the lower limit is preferably 0.0005%. More preferably, it is 0.0010% or more. If Ta exceeds 0.1%, the effect of addition is saturated, quenching hardness is reduced due to excessive carbide formation, and the cost is increased. Therefore, when Ta is contained, the upper limit is 0.1%. It is preferable. It is more preferably 0.05% or less, and even more preferably less than 0.03%.
  • Ni 0.0005% or more and 0.1% or less
  • Ni is an element highly effective in improving toughness and hardenability. If less than 0.0005%, there is no effect of addition, so when Ni is contained, the lower limit is preferably 0.0005%. More preferably, it is 0.0010% or more. If Ni exceeds 0.1%, the effect of addition is saturated and the cost also increases. Therefore, when Ni is contained, the upper limit is preferably made 0.1%. More preferably, it is 0.05% or less.
  • Cu 0.0005% or more and 0.1% or less Cu is an element effective for ensuring hardenability. If less than 0.0005%, the effect of addition is not sufficiently confirmed. Therefore, when Cu is contained, the lower limit is preferably 0.0005%. More preferably, it is 0.0010% or more. If Cu is more than 0.1%, defects during hot rolling tend to occur and the productivity is deteriorated such as a decrease in yield. Therefore, when Cu is contained, the upper limit is preferably 0.1%. More preferably, it is 0.05% or less.
  • V 0.0005% or more and 0.1% or less V, like Nb and Ta, forms carbonitrides to prevent abnormal grain growth of crystal grains during heating before quenching, improve toughness, and improve temper softening resistance. It is an effective element. If it is less than 0.0005%, the effect of addition is not sufficiently exhibited, so when V is contained, the lower limit is preferably made 0.0005%. More preferably, it is 0.0010% or more. If V exceeds 0.1%, not only the effect of addition is saturated, but also the elongation decreases as the tensile strength of the base material increases due to Nb carbide. Therefore, when V is contained, the upper limit is set to 0. It is preferably set to 1%. It is more preferably 0.05% or less, and even more preferably less than 0.03%.
  • W 0.0005% or more and 0.1% or less W, like Nb and V, forms carbonitrides and is effective in preventing abnormal grain growth of austenite crystal grains during heating before quenching and improving temper softening resistance. Is an element. If it is less than 0.0005%, the effect of addition is small, so when W is contained, the lower limit is preferably 0.0005%. More preferably, it is 0.0010% or more. If W exceeds 0.1%, the effect of addition is saturated, the quenching hardness is reduced due to excessive carbide formation, and the cost increases, so the upper limit is made 0.1% when W is contained. It is preferable. It is more preferably 0.05% or less, and even more preferably less than 0.03%.
  • the microstructure has ferrite and cementite
  • the cementite has a circle equivalent diameter of 0.1 ⁇ m or less with respect to the total cementite number of 20% or less, and an average cementite diameter of 2.5 ⁇ m or less
  • the area ratio of the cementite to the total microstructure is 3.5% or more and 10.0% or less
  • the average concentration of the solid solution B in the region from the surface layer to the depth of 100 ⁇ m is 10 mass ppm or more
  • the average concentration of N present as AlN in the region from the surface layer to a depth of 100 ⁇ m is 70 mass ppm or less.
  • the average particle diameter of the ferrite is 4 to 25 ⁇ m. More preferably, it is 5 ⁇ m or more.
  • the microstructure of the high carbon hot-rolled steel sheet of the present invention has ferrite and cementite.
  • the area ratio of ferrite is preferably 90% or more. If the ferrite area ratio is less than 90%, the formability is deteriorated, and cold working may be difficult for parts with high workability. Therefore, the ferrite area ratio is preferably 90% or more. More preferably, it is 92% or more.
  • pearlite may be generated in addition to the above-mentioned ferrite and cementite. If the area ratio of pearlite with respect to the entire microstructure is 6.5% or less, the effect of the present invention is not impaired, and thus it may be included.
  • Ratio of the number of cementites having a circle-equivalent diameter of 0.1 ⁇ m or less to the total number of cementite 20% or less If there is a large amount of cementite having a circle-equivalent diameter of 0.1 ⁇ m or less, it becomes hardened due to dispersion strengthening and elongation is reduced. From the viewpoint of obtaining cold workability, in the present invention, the number of cementites having a circle equivalent diameter of 0.1 ⁇ m or less is 20% or less with respect to the total number of cementites. As a result, it is possible to further achieve a tensile strength of 480 MPa or less and a total elongation (El) of 33% or more.
  • the number of cementites having a circle equivalent diameter of 0.1 ⁇ m or less is preferably 10% or less of the total number of cementites. .. By setting the number of cementites having a circle equivalent diameter of 0.1 ⁇ m or less to 10% or less with respect to the total number of cementites, it is possible to achieve a tensile strength of 440 MPa or less and a total elongation (El) of 36% or more.
  • the reason for defining the proportion of cementite having a circle-equivalent diameter of 0.1 ⁇ m or less is that cementite having a diameter of 0.1 ⁇ m or less produces dispersion strengthening ability, and if the size of cementite increases, cold workability is impaired.
  • the number of cementites having a circle equivalent diameter of 0.1 ⁇ m or less is preferably 3% or more with respect to the total number of cementites.
  • the cementite diameter existing before quenching is about 0.07 to 3.0 ⁇ m in equivalent circle diameter. Therefore, the dispersed state of cementite having a circle equivalent diameter of more than 0.1 ⁇ m before quenching, which is a size that does not significantly affect precipitation strengthening, is not particularly specified in the present invention.
  • Average cementite diameter 2.5 ⁇ m or less
  • the average cementite diameter is set to 2.5 ⁇ m or less. It is more preferably 2.0 ⁇ m or less. If the cementite is too fine, the precipitation strengthening of the cementite deteriorates the cold workability. Therefore, the average cementite diameter is preferably 0.1 ⁇ m or more. More preferably, it is 0.15 ⁇ m or more.
  • cementite diameter refers to a circle-equivalent diameter of cementite
  • the circle-equivalent diameter of cementite is a value obtained by measuring the major axis and the minor axis of cementite and converting them to the circle-equivalent diameters.
  • the “average cementite diameter” refers to a value obtained by dividing the sum of the equivalent circle diameters of all the cementites converted into equivalent circle diameters by the total number of cementites.
  • Proportion (area ratio) of cementite to all microstructures 3.5% or more and 10.0% or less
  • the ratio of cementite to all microstructures exceeds 10.0%, it contributes to precipitation strengthening.
  • the number of cementite particles having a particle size of 0.1 ⁇ m or less increases, and the steel hardens, so the content is made 10.0% or less. It is preferably 9.5% or less.
  • the above ratio is less than 3.5%, the substantial C content does not reach 0.20%, and the predetermined hardness cannot be obtained after heat treatment, so the content is made 3.5% or more. More preferably, it is 4.0% or more.
  • the average ferrite grain size is preferably 25 ⁇ m or less. It is more preferably at least 5 ⁇ m, and even more preferably at least 6 ⁇ m. More preferably, it is 20 ⁇ m or less. More preferably, it is 18 ⁇ m or less.
  • the equivalent circle diameter of cementite, the average cementite diameter, the ratio of cementite to the total microstructure, the area ratio of ferrite, the average particle diameter of ferrite, etc. are measured by the method described in Examples described later. can do.
  • the average concentration of the solute B is 12 mass ppm or more. More preferably, it is 15 mass ppm or more. If the solid solution B is too high, the development of the texture of the hot rolled structure is hindered, so the content is set to 40 mass ppm or less. More preferably, it is 35 mass ppm or less.
  • the average concentration of N amount existing as AlN in the region from the surface layer to the depth of 100 ⁇ m 70 mass ppm or less
  • the average concentration of N amount present as AlN in the region from the steel plate surface layer to the 100 ⁇ m position in the plate thickness direction When the content is 70 mass ppm or less, the growth of crystal grains is promoted in the austenite region in the heating before quenching. This makes it difficult to obtain a structure called pearlite or sorbite in the cooling stage, does not cause insufficient quenching, and has a predetermined surface hardness.
  • the average concentration of the amount of N existing as AlN in the region from the surface layer to the depth of 100 ⁇ m is preferably 50 mass ppm or less.
  • the average concentration of the N content is preferably 10 mass ppm or more, and more preferably 20 mass ppm or more.
  • the amount of solid solution B in the steel sheet surface layer and the amount of N present as AlN are closely related to the manufacturing conditions in each step such as heating conditions, winding conditions, and annealing conditions. It turned out to be necessary to optimize. The reason necessary to obtain the amount of solid solution B and the amount of N as AlN in each step will be described later.
  • the high-carbon hot-rolled steel sheet of the present invention is required to have excellent cold workability because it is used to form automobile parts such as gears, transmissions, and seat recliners by cold pressing. Further, it is necessary to increase hardness by quenching treatment to impart wear resistance. Therefore, the high carbon hot-rolled steel sheet of the present invention is excellent by reducing the tensile strength of the steel sheet to a tensile strength (TS) of 480 MPa or less and increasing the elongation to a total elongation (El) of 33% or more. It has both cold workability and excellent hardenability (dip hardenability, carburizing hardenability). More preferably, TS is 460 MPa or less and El is 35% or more.
  • the high-carbon hot-rolled steel sheet of the present invention is made of steel having the above-described composition, and after this material (steel material) is hot-roughly rolled, finish rolling end temperature: Ar 3 transformation point or higher. After finish rolling, the average cooling rate: 20-100°C/sec, cooling to 650-750°C, winding temperature: 500-700°C, cooling to room temperature to obtain a hot rolled steel sheet Manufactured by heating a hot rolled steel sheet at an average heating rate of 15° C./h or more in a temperature range of 450 to 600° C. and annealing at a temperature of less than Ac 1 transformation point for 1.0 h or more. ..
  • ° C.” regarding temperature indicates the temperature on the surface of the steel plate or the surface of the steel material.
  • the manufacturing method of the steel material does not need to be particularly limited.
  • a converter and an electric furnace can be used to produce the high carbon steel of the present invention.
  • High carbon steel melted by a known method such as a converter is made into a slab (steel material) by ingot-bulk rolling or continuous casting.
  • the slab is usually heated and then hot-rolled (hot rough rolling, finish rolling).
  • Finishing rolling end temperature Finish rolling at Ar 3 transformation point or higher If the finishing rolling termination temperature is less than Ar 3 transformation point, coarse ferrite grains are formed after hot rolling and after annealing, and elongation is remarkably reduced. Therefore, the finish rolling end temperature is set to the Ar 3 transformation point or higher.
  • the temperature is preferably (Ar 3 transformation point+20° C.) or higher.
  • the upper limit of the finish rolling finish temperature is not particularly limited, but it is preferably 1000° C. or lower for smooth cooling after finish rolling.
  • the average cooling rate is 20 to 100° C./sec and is cooled to 650 to 750° C.
  • the average cooling rate from 650 to 750° C. greatly affects the size of spheroidized cementite after annealing. If the average cooling rate after finish rolling is less than 20° C./sec, the ferrite structure and the pearlite structure are too large as the pre-annealing structure, so that the predetermined cementite dispersed state and size cannot be obtained after the annealing. Therefore, it is necessary to cool at 20° C./sec or more. It is preferably 25° C./sec or more.
  • Winding temperature 500-700°C
  • the hot rolled steel sheet after finish rolling is wound into a coil shape. If the coiling temperature is too high, the strength of the hot-rolled steel sheet becomes too low, and when coiled into a coil shape, the coil may be deformed by its own weight. Therefore, it is not preferable from the viewpoint of operation. Therefore, the upper limit of the winding temperature is 700°C. The temperature is preferably 690°C or lower. On the other hand, if the winding temperature is too low, the hot-rolled steel sheet becomes hard, which is not preferable. Therefore, the winding temperature is 500°C. It is preferably 530° C. or higher.
  • Average heating rate in the temperature range of 450 to 600° C. 15° C./h or more
  • the hot rolled steel sheet obtained as described above is annealed (cementite spheroidizing annealing).
  • ammonia gas is likely to be generated in the temperature range of 450 to 600° C., and nitrogen decomposed from the ammonia gas enters the surface steel sheet and combines with B and Al in the steel to form a nitride. To do. Therefore, the heating time in the temperature range of 450 to 600° C. should be as short as possible.
  • the average heating rate in this temperature range is 15° C./h or more. From the viewpoint of suppressing variations in the furnace for the purpose of improving productivity, it is preferably 100° C./h or less, more preferably 70° C./h or less.
  • the following two-step annealing can be applied instead of the above-mentioned annealing.
  • the hot-rolled steel sheet is heated in a temperature range of 450 to 600° C. at an average heating rate of 15° C./h or more to obtain an Ac 1 transformation point. Hold at 0.5 h or more below the Ac 3 transformation point (first annealing), then cool to less than Ar 1 transformation point at average cooling rate: 1 to 20° C./h, and 20 h or more below Ar 1 transformation point It is also possible to manufacture by performing a two-step annealing that holds (second-step annealing).
  • the hot-rolled steel sheet is heated in the temperature range of 450 to 600° C. at an average heating rate of 15° C./h or more and kept at the Ac 1 transformation point or more for 0.5 h or more to precipitate in the hot-rolled steel sheet.
  • the relatively fine carbide is dissolved to form a solid solution in the ⁇ phase, and thereafter, the average cooling rate is cooled to below Ar 1 transformation point at an average cooling rate of 1 to 20° C./h, and maintained for 20 hours or more below Ar 1 transformation point. ..
  • a solid solution C is precipitated by using a relatively coarse undissolved carbide as a nucleus, and the ratio of the number of cementites having a circle-equivalent diameter of 0.1 ⁇ m or less to the total number of cementites is 20% or less ( Cementite) dispersion can be controlled. That is, in the present invention, the two-step annealing is performed under predetermined conditions to control the dispersed form of the carbide and soften the steel sheet. In the high carbon steel sheet targeted by the present invention, it is important to control the dispersed form of carbides after annealing in order to soften the steel.
  • first-stage annealing by holding the high carbon hot-rolled steel sheet at the Ac 1 transformation point or more and the Ac 3 transformation point or less (first-stage annealing), fine carbides are dissolved and C is solidified in ⁇ (austenite). Melt.
  • second annealing the ⁇ / ⁇ interface and undissolved carbides existing in the temperature range above the Ac 1 transformation point become nucleation sites and are relatively coarse. Carbide precipitates.
  • the atmosphere gas at the time of annealing any of nitrogen, hydrogen, and a mixed gas of nitrogen and hydrogen can be used.
  • Average heating rate in the temperature range of 450 to 600° C. 15° C./h or more
  • ammonia gas is easily generated in the temperature range of 450 to 600° C., and nitrogen decomposed from the ammonia gas becomes surface steel sheet.
  • the heating time in the temperature range of 450 to 600° C. is made as short as possible because it enters and combines with B and Al in the steel to form a nitride.
  • the average heating rate in this temperature range is 15° C./h or more. It is preferably 20° C./h or more.
  • the upper limit of the average heating rate is preferably 100°C/h, more preferably 90°C/h or less.
  • the annealing temperature of the first step exceeds the Ac 3 transformation point, a large number of rod-shaped cementites are obtained after annealing and a predetermined elongation cannot be obtained, so the temperature is set to the Ac 3 transformation point or lower. Further, in the present invention, if the holding time at the Ac 1 transformation point or more and the Ac 3 transformation point or less is less than 0.5 h, fine carbides cannot be sufficiently dissolved. For this reason, as the first-stage annealing, 0.5 h or more is maintained at the Ac 1 transformation point or more and the Ac 3 transformation point or less.
  • the holding time is preferably 1.0 h or longer.
  • the holding time is preferably 10 hours or less. Even when annealing is performed while maintaining the temperature at the Ac 1 transformation point or more and the Ac 3 transformation point or less, the heating rate is 15°C/h or more in the average heating rate in the temperature range of 450 to 600°C, and the upper limit is 100°C/ It is preferably h or less.
  • Average cooling rate cooling to below Ar 1 transformation point at 1 to 20° C./h
  • average cooling rate 1 Cool at ⁇ 20°C/h.
  • C discharged from the austenite along with the transformation from austenite to ferrite is precipitated as a relatively coarse spherical carbide by using the ⁇ / ⁇ interface and undissolved carbide as a nucleation site. In this cooling, it is necessary to adjust the cooling rate so that pearlite is not generated.
  • the average cooling rate from the first annealing to the second annealing is less than 1° C./h, the production efficiency is poor, so the average cooling rate is 1° C./h or more. It is preferably 5° C./h or more.
  • the rate is set to 20° C./h or less. The rate is preferably 15° C./h or less.
  • Second-stage annealing After the annealing in the first step described above, cooling is performed at a predetermined average cooling rate and the temperature is maintained below the Ar 1 transformation point, whereby coarse spherical carbides are further grown and fine carbides disappear by Ostwald ripening. If the holding time below the Ar 1 transformation point is less than 20 h, the carbide cannot be grown sufficiently and the hardness after annealing becomes too large. Therefore, the second annealing is held for 20 hours or more below the Ar 1 transformation point.
  • the second annealing temperature is preferably 660° C. or higher in order to sufficiently grow the carbide, and the holding time is 30 h or less from the viewpoint of production efficiency. preferable.
  • the above Ac 3 transformation point, Ac 1 transformation point, Ar 3 transformation point, and Ar 1 transformation point may be determined by actual measurement by thermal expansion measurement or electric resistance measurement during heating or cooling by the Formaster test or the like. it can.
  • the above average heating rate and average cooling rate are obtained by measuring the temperature with a thermocouple installed in the furnace.
  • test pieces were sampled from the hot rolled annealed sheet thus obtained, and the microstructure, the amount of solid solution B, the amount of N in AlN, the tensile strength, the total elongation and Hardening hardness (steel plate hardness after quenching, steel plate hardness after carburizing and quenching) was determined.
  • the Ac 3 transformation point, the Ac 1 transformation point, the Ar 1 transformation point and the Ar 3 transformation point shown in Table 1 were obtained by the Formaster test.
  • the area ratio (%) of the ferrite was obtained by binarizing the ferrite and the area other than the ferrite using image analysis software from the SEM image.
  • the area ratio (%) of cementite was obtained by binarizing the cementite and the region other than the cementite from the SEM image using image analysis software.
  • the value obtained by subtracting the area ratio (%) of each of ferrite and cementite from 100 (%) was defined as the area ratio (%) of pearlite.
  • individual cementite diameters were evaluated for the taken micrographs.
  • the cementite diameter the major axis and the minor axis were measured and converted into a circle equivalent diameter.
  • the average cementite diameter was calculated by dividing the sum of the equivalent circle diameters of all the cementites converted to the equivalent circle diameter by the total number of cementites.
  • the number of cementites having a circle equivalent diameter of 0.1 ⁇ m or less was measured and used as the number of cementite having a circle equivalent diameter of 0.1 ⁇ m or less.
  • the total number of cementites was calculated and used as the total number of cementites.
  • the ratio of the number of cementites having a circle-equivalent diameter of 0.1 ⁇ m or less to the total number of cementites ((the number of cementites having a circle-equivalent diameter of 0.1 ⁇ m or less/total number of cementites) ⁇ 100(%)) was determined.
  • the "ratio of cementite having a circle-equivalent diameter of 0.1 ⁇ m or less” may be simply referred to as cementite having a circle-equivalent diameter of 0.1 ⁇ m or less.
  • the average grain size of ferrite was determined for the photographed structure using the grain size evaluation method (cutting method) specified in JIS G 0551.
  • the quenching hardness is the hardness of the cut surface of the test piece after quenching under the condition of a load of 0.2 kgf with a Vickers hardness tester in an area within the thickness of 70 ⁇ m from the surface layer and a quarter thickness. Five points were measured and the average hardness was determined, which was taken as the quenching hardness (HV). The region within the plate thickness of 70 ⁇ m from the surface layer is indicated as “surface layer” in Table 2-2 and Table 3-2.
  • Table 4 shows the acceptance criteria of the hardenability according to the C content, which can be evaluated as having sufficient hardenability.
  • the ratio of the number of cementites having a circle equivalent diameter of 0.1 ⁇ m or less to the total number of cementites was 20% or less, and the average cementite diameter was Is 2.5 ⁇ m or less, the ratio of the cementite to the total microstructure is 3.5% or more and 10.0% or less, and has a microstructure having ferrite and cementite, and is excellent in cold workability and hardenability. It turns out that it is also excellent. Also, excellent mechanical properties such as tensile strength of 480 MPa or less and total elongation (El) of 33% or more could be obtained.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Heat Treatment Of Sheet Steel (AREA)

Abstract

Provided are a high-carbon hot-rolled steel sheet and a method for manufacturing the same. The present invention is a high-carbon hot-rolled steel sheet having a specific component composition, wherein: the microstructure comprises ferrite, cementite, and pearlite accounting for 6.5% or less, in area percentage, of the whole microstructure; as regards the cementite, the proportion of the quantity of cementite having a circle-equivalent diameter of 0.1 µm or less is 20% or less relative to the total quantity of cementite, the average cementite diameter is 2.5 µm or less, and the proportion of cementite relative to the whole microstructure is from 3.5% to 10.0% in area percentage; the average concentration of the amount of solid-soluted B in a region up to a depth of 100 µm from the surface layer is 10 ppm by mass or greater; and the average concentration of the amount of N existing as AlN in a region up to a depth of 100 µm from the surface layer is 70 ppm by mass or less.

Description

高炭素熱延鋼板およびその製造方法High carbon hot rolled steel sheet and method for producing the same
　本発明は、冷間加工性および焼入れ性（ズブ焼入れ性および浸炭焼入れ性）に優れる高炭素熱延鋼板およびその製造方法に関する。 The present invention relates to a high carbon hot-rolled steel sheet excellent in cold workability and hardenability (dip hardenability and carburizing hardenability) and a method for producing the same.
　現在、トランスミッション、シートリクライナーなどの自動車用部品は、ＪＩＳ　Ｇ４０５１に規定された機械構造用炭素鋼鋼材および機械構造用合金鋼鋼材である熱延鋼板（高炭素熱延鋼板）を、冷間加工によって所望の形状に加工した後、所望の硬さを確保するために焼入れ処理を施して製造されることが多い。このため、素材となる熱延鋼板には優れた冷間加工性や焼入れ性が必要とされ、これまでに種々の鋼板が提案されている。 Currently, automotive parts such as transmissions and seat recliners are produced by cold working hot-rolled steel sheets (high carbon hot-rolled steel sheets), which are carbon steel steels for machine structures and alloy steels for machine structures specified in JIS G4051. After being processed into a desired shape, it is often manufactured by quenching to secure a desired hardness. Therefore, the hot-rolled steel sheet used as a material is required to have excellent cold workability and hardenability, and various steel sheets have been proposed so far.
　例えば、特許文献１には、重量％で、Ｃ：０．１５～０．９％、Ｓｉ：０．４％以下、Ｍｎ：０．３～１．０％、Ｐ：０．０３％以下、Ｔ．Ａｌ：０．１０％以下、さらにＣｒ：１．２％以下、Ｍｏ：０．３％以下、Ｃｕ：０．３％以下、Ｎｉ：２．０％以下のうち１種以上あるいはＴｉ：０．０１～０．０５％、Ｂ：０．０００５～０．００５％、Ｎ：０．０１％以下を含有する成分組成とし、球状化率８０％以上、平均粒径０．４～１．０μｍの炭化物がフェライト中に分散した組織をもつ精密打抜き用高炭素鋼板が記載されている。 For example, in Patent Document 1, C: 0.15 to 0.9%, Si: 0.4% or less, Mn: 0.3 to 1.0%, P: 0.03% or less in weight%. T. Al: 0.10% or less, Cr: 1.2% or less, Mo: 0.3% or less, Cu: 0.3% or less, Ni: 2.0% or less, or Ti: 0. 01 to 0.05%, B: 0.0005 to 0.005%, N: 0.01% or less, a spheroidization rate of 80% or more, and an average particle size of 0.4 to 1.0 μm. A high carbon steel sheet for precision punching is described which has a structure in which carbides are dispersed in ferrite.
　特許文献２には、質量％でＣ：０．２％以上、Ｔｉ：０．０１～０．０５％、Ｂ：０．０００３～０．００５％を含有する成分組成とし、炭化物の平均粒径が１．０μｍ以下、かつ０．３μｍ以下の炭化物の比率が２０％以下である加工性を改善した高炭素鋼板が記載されている。 In Patent Document 2, the composition is such that C: 0.2% or more, Ti: 0.01 to 0.05%, and B: 0.0003 to 0.005% by mass%, and the average particle diameter of carbide is Describes a high-carbon steel sheet having improved workability in which the ratio of carbides having a grain size of 1.0 μm or less and 0.3 μm or less is 20% or less.
　特許文献３には、質量％で、Ｃ：０．２０％以上０．４５％以下、Ｓｉ：０．０５％以上０．８％以下、Ｍｎ：０．５％以上２．０％以下、Ｐ：０．００１％以上０．０４％以下、Ｓ：０．０００１％以上０．００６％以下、Ａｌ：０．００５％以上０．１％以下、Ｔｉ：０．００５％以上０．２％以下、Ｂ：０．００１％以上０．０１％以下、及びＮ：０．０００１％以上０．０１％以下、さらにＣｒ：０．０５％以上０．３５％以下、Ｎｉ：０．０１％以上１．０％以下、Ｃｕ：０．０５％以上０．５％以下、Ｍｏ：０．０１％以上１．０％以下、Ｎｂ：０．０１％以上０．５％以下、Ｖ：０．０１％以上０．５％以下、Ｔａ：０．０１％以上０．５％以下、Ｗ：０．０１％以上０．５％以下、Ｓｎ：０．００３％以上０．０３％以下、Ｓｂ：０．００３％以上０．０３％以下、Ａｓ：０．００３％以上０．０３％以下の１種または２種以上の成分を有するＢ添加鋼が記載されている。 In Patent Document 3, C: 0.20% or more and 0.45% or less, Si: 0.05% or more and 0.8% or less, Mn: 0.5% or more and 2.0% or less, P in mass% : 0.001% to 0.04%, S: 0.0001% to 0.006%, Al: 0.005% to 0.1%, Ti: 0.005% to 0.2% , B: 0.001% or more and 0.01% or less, and N: 0.0001% or more and 0.01% or less, Cr: 0.05% or more and 0.35% or less, Ni: 0.01% or more 1 0.0% or less, Cu: 0.05% or more and 0.5% or less, Mo: 0.01% or more and 1.0% or less, Nb: 0.01% or more and 0.5% or less, V: 0.01% Or more and 0.5% or less, Ta: 0.01% or more and 0.5% or less, W: 0.01% or more and 0.5% or less, Sn: 0.003% or more and 0.03% or less, Sb: 0. A B-added steel having one or more components of 003% to 0.03% and As: 0.003% to 0.03% is described.
　特許文献４には、質量％で、Ｃ：０．１０～１．２％、Ｓｉ：０．０１～２．５％、Ｍｎ：０．１～１．５％、Ｐ：０．０４％以下、Ｓ：０．０００５～０．０５％、Ａｌ：０．２％以下、Ｔｅ：０．０００５～０．０５％、Ｎ：０．０００５～０．０３％、さらにＳｂ：０．００１～０．０５％、加えてＣｒ：０．２～２．０％、Ｍｏ：０．１～１．０％、Ｎｉ：０．３～１．５％、Ｃｕ：１．０％以下、Ｂ：０．００５％以下のうち１種以上を含有する成分組成とし、フェライトとパーライトを主体とする組織からなり、フェライト結晶粒度が１１番以上である冷間加工性と低脱炭性を改善した機械構造用鋼が記載されている。 In Patent Document 4, C: 0.10 to 1.2%, Si: 0.01 to 2.5%, Mn: 0.1 to 1.5%, and P: 0.04% or less in mass%. , S: 0.0005 to 0.05%, Al: 0.2% or less, Te: 0.0005 to 0.05%, N: 0.0005 to 0.03%, and Sb: 0.001 to 0 0.05%, Cr: 0.2-2.0%, Mo: 0.1-1.0%, Ni: 0.3-1.5%, Cu: 1.0% or less, B:0 A mechanical structure with a composition containing at least one of 0.005% or less, a structure mainly composed of ferrite and pearlite, and having a ferrite grain size of 11 or more and improved cold workability and low decarburization. Steel for use is described.
　特許文献５には、質量％で、Ｃ：０．２０～０．４０％、Ｓｉ：０．１０％以下、Ｍｎ：０．５０％以下、Ｐ：０．０３％以下、Ｓ：０．０１０％以下、ｓｏｌ．Ａｌ：０．１０％以下、Ｎ：０．００５％以下、Ｂ：０．０００５～０．００５０％を含有し、さらにＳｂ、Ｓｎ、Ｂｉ、Ｇｅ、Ｔｅ、Ｓｅのうち１種以上を合計で０．００２～０．０３％含有し、フェライトとセメンタイトからなり、フェライト粒内のセメンタイト密度が０．１０個／μｍ２以下であるミクロ組織を有し、硬さがＨＲＢで７５以下、全伸びが３８％以上である焼入れ性および加工性を改善した高炭素熱延鋼板が記載されている。 In Patent Document 5, C: 0.20 to 0.40%, Si: 0.10% or less, Mn: 0.50% or less, P: 0.03% or less, and S: 0.010 in mass%. % Or less, sol. Al: 0.10% or less, N: 0.005% or less, B: 0.0005 to 0.0050%, and one or more of Sb, Sn, Bi, Ge, Te and Se in total. It contains 0.002 to 0.03%, consists of ferrite and cementite, has a microstructure with a cementite density of 0.10 particles/μm 2 or less in ferrite grains, and has a hardness of 75 or less in HRB and a total elongation. Is 38% or more, and a high carbon hot rolled steel sheet having improved hardenability and workability is described.
　特許文献６には、質量％で、Ｃ：０．２０～０．４８％、Ｓｉ：０．１０％以下、Ｍｎ：０．５０％以下、Ｐ：０．０３％以下、Ｓ：０．０１０％以下、ｓｏｌ．Ａｌ：０．１０％以下、Ｎ：０．００５％以下、Ｂ：０．０００５～０．００５０％を含有し、さらにＳｂ、Ｓｎ、Ｂｉ、Ｇｅ、Ｔｅ、Ｓｅのうち１種以上を合計で０．００２～０．０３％含有し、フェライトとセメンタイトからなり、前記フェライト粒内のセメンタイト密度が０．１０個/μｍ２以下であるミクロ組織を有し、硬さがＨＲＢで６５以下、全伸びが４０％以上である焼入れ性および加工性を改善した高炭素熱延鋼板が記載されている。 In Patent Document 6, in mass%, C: 0.20 to 0.48%, Si: 0.10% or less, Mn: 0.50% or less, P: 0.03% or less, S: 0.010. % Or less, sol. Al: 0.10% or less, N: 0.005% or less, B: 0.0005 to 0.0050%, and one or more of Sb, Sn, Bi, Ge, Te and Se in total. It contains 0.002 to 0.03%, is composed of ferrite and cementite, has a microstructure with a cementite density of 0.10 particles/μm 2 or less in the ferrite grains, and has a hardness of HRB of 65 or less, and a total hardness of 65 or less. A high carbon hot rolled steel sheet having an elongation of 40% or more and improved hardenability and workability is described.
　特許文献７には、質量％で、Ｃ：０．２０～０．４０％、Ｓｉ：０．１０％以下、Ｍｎ：０．５０％以下、Ｐ：０．０３％以下、Ｓ：０．０１０％以下、ｓｏｌ．Ａｌ：０．１０％以下、Ｎ：０．００５％以下、Ｂ：０．０００５～０．００５０％を含有し、さらにＳｂ、Ｓｎ、Ｂｉ、Ｇｅ、Ｔｅ、Ｓｅのうち１種以上を合計で０．００２～０．０３％含有し、Ｂ含有量に占める固溶Ｂ量の割合が７０％以上であり、フェライトとセメンタイトからなり、フェライト粒内のセメンタイト密度が０．０８個／μｍ２以下であるミクロ組織を有し、硬さがＨＲＢで７３以下、全伸びが３９％以上である高炭素熱延鋼板が記載されている。 In Patent Document 7, C: 0.20 to 0.40%, Si: 0.10% or less, Mn: 0.50% or less, P: 0.03% or less, and S: 0.010 in mass%. % Or less, sol. Al: 0.10% or less, N: 0.005% or less, B: 0.0005 to 0.0050%, and one or more of Sb, Sn, Bi, Ge, Te and Se in total. The content of 0.002 to 0.03%, the proportion of solid solution B in the B content is 70% or more, and it is composed of ferrite and cementite, and the cementite density in the ferrite grains is 0.08 pieces/μm 2 or less. And a high carbon hot rolled steel sheet having a HRB of 73 or less and a total elongation of 39% or more.
　特許文献８には、質量％で、Ｃ：０．１５～０．３７％、Ｓｉ：１％以下、Ｍｎ：２．５％以下、Ｐ：０．１％以下、Ｓ：０．０３％以下、ｓｏｌ．Ａｌ：０．１０％以下、Ｎ：０．０００５～０．００５０％、Ｂ：０．００１０～０．００５０％、およびＳｂ、Ｓｎのうち少なくとも１種：合計で０．００３～０．１０％を含有し、かつ０．５０≦（１４［Ｂ］）／（１０．８［Ｎ］）の関係を満足し、残部がＦｅおよび不可避的不純物からなる組成を有し、フェライト相とセメンタイトからなり、フェライト相の平均粒径が１０μｍ以下、セメンタイトの球状化率が９０％以上であるミクロ組織を有し、全伸びが３７％以上ある高炭素熱延鋼板が記載されている。 In Patent Document 8, C: 0.15 to 0.37%, Si: 1% or less, Mn: 2.5% or less, P: 0.1% or less, S: 0.03% or less in mass%. , Sol. Al: 0.10% or less, N: 0.0005 to 0.0050%, B: 0.0010 to 0.0050%, and at least one of Sb and Sn: 0.003 to 0.10% in total And satisfying the relationship of 0.50≦(14[B])/(10.8[N]), the balance being Fe and inevitable impurities, and consisting of a ferrite phase and cementite. A high carbon hot-rolled steel sheet having a microstructure in which the average grain size of the ferrite phase is 10 μm or less, the spheroidization rate of cementite is 90% or more, and the total elongation is 37% or more is described.
特開２００９－２９９１８９号公報JP, 2009-299189, A 特開２００５－３４４１９４号公報JP 2005-344194 A 特許第４０１２４７５号公報Japanese Patent No. 4012475 特許第４７８２２４３号公報Japanese Patent No. 4782243 特開２０１５－０１７２８３号公報JP, 2005-017283, A 特開２０１５－０１７２８４号公報JP, 2005-017284, A 国際公開第２０１５/１４６１７３号International Publication No. 2015/146173 特許第５４５８６４９号公報Japanese Patent No. 5458649
　特許文献１に記載される技術は、精密打抜き性に関するものであり、炭化物の分散形態が精密打抜き性及び焼入れ性に及ぼす影響を記載している。具体的には、特許文献１では、平均炭化物粒径を０．４～１．０μｍに制御し、球状化率を８０％以上とすることで、精密打抜き性と焼入れ性を改善する鋼板が得られることを記載している。しかし、特許文献１には冷間加工性に関する議論はなく、また浸炭焼入れ性に関する記載もない。 The technology described in Patent Document 1 relates to precision punchability, and describes the influence of the dispersed form of carbide on the precision punchability and hardenability. Specifically, in Patent Document 1, by controlling the average carbide grain size to 0.4 to 1.0 μm and setting the spheroidization rate to 80% or more, a steel sheet that improves precision punchability and hardenability is obtained. It is described that it is possible. However, Patent Document 1 does not discuss cold workability and does not describe carburizing and quenching properties.
　特許文献２に記載される技術は、炭化物平均粒径だけでなく、０．３μｍ以下の微細炭化物が加工性に影響することに注目し、炭化物の平均粒径を１．０μｍ以下に制御し、加えて０．３μｍ以下の炭化物割合を２０％以下に制御する。これにより、加工性を改善した鋼板が得られることを記載しており、さらにＴｉやＢを添加した焼入れ性に優れた鋼板を記載している。しかし、特許文献２では、焼入れ性に影響する固溶Ｂ等の記述はなく、鋼板のどの位置における焼入れ硬さに相当するかについても記述されていない。 The technique described in Patent Document 2 pays attention not only to the average grain size of carbides, but also to the fact that fine carbides of 0.3 μm or less affect workability, and controls the average grain size of carbides to 1.0 μm or less, In addition, the proportion of carbides of 0.3 μm or less is controlled to 20% or less. This describes that a steel sheet with improved workability can be obtained, and further, a steel sheet having Ti and B added and having excellent hardenability is described. However, Patent Document 2 does not describe solid solution B or the like that affects the hardenability, and does not describe at which position of the steel plate the hardened hardness corresponds.
　特許文献３に記載される技術は、成分組成を調整することで、冷間加工性と耐脱炭性を改善した鋼が得られることを記載している。しかしながら、特許文献３には、ズブ焼入れ性、浸炭焼入れ性に関する記載はない。 The technology described in Patent Document 3 describes that a steel having improved cold workability and decarburization resistance can be obtained by adjusting the composition of components. However, in Patent Document 3, there is no description regarding the dip quenching property and the carburizing quenching property.
　特許文献４に記載される技術は、Ｂ、さらにＣｒ、Ｎｉ、Ｃｕ、Ｍｏ、Ｎｂ、Ｖ、Ｔａ、Ｗ、Ｓｎ、Ｓｂ、Ａｓの１種または２種以上の成分を含有し、表層における固溶Ｂを所定量確保することで高い焼入れ性を達成する鋼が得られることを述べている。しかし、特許文献４では、焼鈍工程における雰囲気中の水素濃度が９５％以上と規定されており、窒素雰囲気の焼鈍工程において吸窒を抑えて固溶Ｂを確保することが可能かに関する記載はない。 The technique described in Patent Document 4 contains B and further one or more components of Cr, Ni, Cu, Mo, Nb, V, Ta, W, Sn, Sb, As, and a solid layer in the surface layer. It is stated that by ensuring a predetermined amount of molten B, a steel that achieves high hardenability can be obtained. However, in Patent Document 4, the hydrogen concentration in the atmosphere in the annealing step is specified to be 95% or more, and there is no description on whether it is possible to suppress nitrification and secure the solid solution B in the annealing step in the nitrogen atmosphere. ..
　特許文献５～７に記載される技術は、Ｂ、さらにＳｂ、Ｓｎ、Ｂｉ、Ｇｅ、Ｔｅ、Ｓｅのうち１種以上を合計で０．００２～０．０３％含有することで浸窒防止効果が高く、例えば窒素雰囲気で焼鈍した場合においても、浸窒を防止し、固溶Ｂが所定量維持されることで焼入れ性を高くすることが記載されている。しかしながら、特許文献５～７には、いずれも表層における焼入れ硬さに関する記述はない。 The technologies described in Patent Documents 5 to 7 have the effect of preventing nitriding by containing 0.002 to 0.03% of B and at least one of Sb, Sn, Bi, Ge, Te and Se in total. It is described that, even when annealing is performed in a nitrogen atmosphere, for example, nitrification is prevented and the solid solution B is maintained at a predetermined amount to enhance the hardenability. However, in any of Patent Documents 5 to 7, there is no description about quenching hardness in the surface layer.
　特許文献８に記載される技術では、Ｃ：０．１５～０．３７％でＢとＳｂ、Ｓｎの１種以上を含有することで焼入れ性の高い鋼を提案している。しかしながら、特許文献８では、浸炭焼入れ性といった、より高い焼入れ性については検討されていない。 The technology described in Patent Document 8 proposes a steel with high hardenability by containing C: 0.15 to 0.37% and at least one of B, Sb, and Sn. However, Patent Document 8 does not consider higher quenchability such as carburizing quenchability.
　本発明は、上記問題に鑑みてなされたものであり、優れた冷間加工性および優れた焼入れ性（ズブ焼入れ性、浸炭焼入れ性）を有する高炭素熱延鋼板およびその製造方法を提供することを目的とする。 The present invention has been made in view of the above problems, and provides a high-carbon hot-rolled steel sheet having excellent cold workability and excellent hardenability (dub hardenability, carburizing hardenability) and a method for producing the same. With the goal.
　本発明者らは、上記課題を達成するため、鋼の成分組成として、Ｂ、さらにＳｎおよびＳｂから選んだ１種または２種を含有した高炭素熱延鋼板の製造条件と、冷間加工性および焼入れ性（ズブ焼入れ性、浸炭焼入れ性）との関係について鋭意検討した。その結果、以下の知見を得た。 MEANS TO SOLVE THE PROBLEM In order to achieve the said subject, these artificers have produced the manufacturing conditions of the high carbon hot-rolled steel sheet containing B, and 1 type or 2 types further selected from Sn and Sb as a component composition of steel, and cold workability. And the relationship with the hardenability (dip hardenability, carburizing hardenability) was earnestly examined. As a result, the following findings were obtained.
　ｉ）窒素雰囲気で焼鈍を施す場合、雰囲気中の窒素が浸窒して鋼板中に濃化し、鋼板中のＢやＡｌと結合して表層にＢ窒化物およびＡｌ窒化物を生成する。これにより、鋼板中の固溶Ｂ量が低下すること、あるいはＡｌ窒化物の存在により焼入れ前のオーステナイト域での加熱中にオーステナイト粒径が小さくなることで、焼入れ不足になる場合がある。そのため、本発明では、窒素雰囲気で焼鈍を施す場合、より高い焼入れ性（高い浸炭焼入れ性）が求められる鋼板に対して、ＳｂとＳｎの少なくとも１種以上を鋼中に所定量添加する。また、焼鈍において４５０～６００℃の温度範囲を所定の加熱速度で加熱することで、雰囲気から鋼中への浸窒を所定量に抑制することが可能である。これらにより、上述の浸窒を防止し、固溶Ｂ量の低下およびＡｌ窒化物の増加を抑制することで、より高い焼入れ性（高い浸炭焼入れ性）を確保することが可能である。 I) When annealing is performed in a nitrogen atmosphere, nitrogen in the atmosphere is nitrided and concentrated in the steel sheet, and is combined with B and Al in the steel sheet to form a B nitride and an Al nitride in the surface layer. As a result, the amount of solid solution B in the steel sheet decreases, or the austenite grain size decreases during heating in the austenite region before quenching due to the presence of Al nitride, which may result in insufficient quenching. Therefore, in the present invention, when annealing is performed in a nitrogen atmosphere, a predetermined amount of at least one of Sb and Sn is added to the steel for a steel plate required to have higher hardenability (high carburizing and quenching property). Further, by heating the temperature range of 450 to 600° C. at a predetermined heating rate during annealing, it is possible to suppress nitrification from the atmosphere into the steel to a predetermined amount. With these, it is possible to secure higher quenchability (high carburizing quenchability) by preventing the above-mentioned nitriding and suppressing the decrease of the amount of solid solution B and the increase of Al nitride.
　ｉｉ）冷間加工性、焼入れ前の高炭素熱延鋼板における硬度（硬さ）、全伸び（以下、単に伸びと称する場合もある。）には、円相当直径が０．１μｍ以下のセメンタイトが大きく影響している。円相当直径が０．１μｍ以下のセメンタイト数を全セメンタイト数に対して２０％以下とすることで、引張強度４８０ＭＰａ以下、全伸び（Ｅｌ）が３３％以上を得ることができる。 ii) For cold workability, hardness (hardness) in high carbon hot rolled steel sheet before quenching, and total elongation (hereinafter sometimes simply referred to as elongation), cementite having an equivalent circle diameter of 0.1 μm or less is It has a great influence. By setting the number of cementites having a circle-equivalent diameter of 0.1 μm or less to 20% or less with respect to the total number of cementites, tensile strength of 480 MPa or less and total elongation (El) of 33% or more can be obtained.
　ｉｉｉ) 焼入れ前の高炭素熱延鋼板における硬度（硬さ）、全伸びには、円相当直径が０．１μｍ以下のセメンタイトが大きく影響している。円相当直径が０．１μｍ以下のセメンタイト数を全セメンタイト数に対して１０％以下とすることで、引張強度４４０ＭＰａ以下、全伸び（Ｅｌ）が３６％以上を得ることができる。 Iii) Cementite with a circle equivalent diameter of 0.1 μm or less greatly affects the hardness (hardness) and total elongation of the high carbon hot rolled steel sheet before quenching. By setting the number of cementites having a circle-equivalent diameter of 0.1 μm or less to 10% or less with respect to the total number of cementites, tensile strength of 440 MPa or less and total elongation (El) of 36% or more can be obtained.
　ｉｖ）熱間粗圧延後、仕上圧延終了温度：Ａｒ３変態点以上で仕上圧延を行い、その後平均冷却速度：２０～１００℃／ｓｅｃで６５０～７５０℃まで冷却し、巻取温度：５００～７００℃で巻き取り、常温まで冷却し、熱延鋼板とした後、該熱延鋼板を、平均加熱速度：１５℃／ｈ以上で４５０～６００℃間を加熱し、焼鈍温度：Ａｃ１変態点未満で１．０ｈ以上保持する焼鈍により、所定のミクロ組織を確保できる。 iv) After hot rough rolling, finish rolling is performed at a finish rolling finish temperature: Ar 3 transformation point or higher, and then cooled to 650 to 750° C. at an average cooling rate of 20 to 100° C./sec, and a winding temperature: 500 to After winding at 700° C., cooling to room temperature to form a hot-rolled steel sheet, the hot-rolled steel sheet is heated at an average heating rate of 15° C./h or more between 450 to 600° C., and an annealing temperature: Ac 1 transformation point. Predetermined microstructure can be secured by annealing for 1.0 h or more.
　ｖ）あるいは、熱間粗圧延後、仕上圧延終了温度：Ａｒ３変態点以上の仕上げ圧延を行い、その後平均冷却速度：２０～１００℃／ｓｅｃで６５０～７５０℃まで冷却し、巻取温度：５００～７００℃で巻き取り、常温まで冷却し、熱延鋼板とした後、該熱延鋼板を、平均加熱速度：１５℃／ｈ以上で４５０～６００℃間を加熱し、Ａｃ１変態点以上Ａｃ３変態点以下で０．５ｈ以上保持し、次いで平均冷却速度：１～２０℃／ｈでＡｒ１変態点未満に冷却して、Ａｒ１変態点未満で２０ｈ以上保持するといった２段焼鈍により、所定のミクロ組織を確保できる。 v) Alternatively, after hot rough rolling, finish rolling is performed: finish rolling at a finish temperature of Ar 3 transformation point or higher, and thereafter, cooling is performed at an average cooling rate of 20 to 100° C./sec to 650 to 750° C., and a winding temperature is: After winding at 500 to 700°C and cooling to room temperature to form a hot rolled steel sheet, the hot rolled steel sheet is heated between 450 and 600°C at an average heating rate of 15°C/h or more to obtain an Ac 1 transformation point or more. By a two-stage annealing, in which the temperature is kept below the Ac 3 transformation point for 0.5 h or more, then the average cooling rate is cooled to below the Ar 1 transformation point at an average cooling rate of 1 to 20° C./h, and held below the Ar 1 transformation point for 20 h or more. A predetermined microstructure can be secured.
　本発明は以上の知見に基づいてなされたものであり、以下を要旨とするものである。
［１］質量％で、Ｃ：０．２０％以上０．５０％以下、Ｓｉ：０．８％以下、Ｍｎ：０．１０％以上０．８０％以下、Ｐ：０．０３％以下、Ｓ：０．０１０％以下、ｓｏｌ．Ａｌ：０．１０％以下、Ｎ：０．０１％以下、Ｃｒ：１．０％以下、Ｂ：０．０００５％以上０．００５％以下、さらにＳｂおよびＳｎから選んだ１種または２種を合計で０．００２％以上０．１％以下を含有し、残部がＦｅおよび不可避的不純物からなる成分組成を有し、ミクロ組織は、フェライト、セメンタイト、および全ミクロ組織に対して面積率で６.５％以下の割合を占めるパーライトを有し、前記セメンタイトは、全セメンタイト数に対する円相当直径０．１μｍ以下のセメンタイト数の割合が２０％以下であり、平均セメンタイト径が２．５μｍ以下、全ミクロ組織に対する前記セメンタイトの占める割合が面積率で３．５％以上１０．０％以下であり、表層から深さ１００μｍまでの領域における固溶Ｂ量の平均濃度が１０質量ｐｐｍ以上であり、表層から深さ１００μｍまでの領域におけるＡｌＮとして存在するＮ量の平均濃度が７０質量ｐｐｍ以下である高炭素熱延鋼板。
［２］引張強度が４８０ＭＰａ以下、全伸びが３３％以上である［１］に記載の高炭素熱延鋼板。
［３］前記フェライトの平均粒径が４～２５μｍである［１］または［２］に記載の高炭素熱延鋼板。
［４］前記成分組成に加えてさらに、質量％で、下記Ａ群およびＢ群のうちから選ばれた１群または２群を含有する[１]～[３]のいずれかに記載の高炭素熱延鋼板。
　　　　　　　　　　　　　　　　記
Ａ群：Ｔｉ：０．０６％以下
Ｂ群：Ｎｂ、Ｍｏ、Ｔａ、Ｎｉ、Ｃｕ、Ｖ、Ｗのうちから選ばれた１種または２種以上を、それぞれ０．０００５％以上０．１％以下
［５］［１］～［４］のいずれかに記載の高炭素熱延鋼板の製造方法であって、前記成分組成を有する鋼を、熱間粗圧延後、仕上圧延終了温度：Ａｒ３変態点以上で仕上圧延を行い、その後、平均冷却速度：２０～１００℃／ｓｅｃで６５０～７５０℃まで冷却し、巻取温度：５００～７００℃で巻き取り、熱延鋼板とした後、該熱延鋼板を、平均加熱速度：１５℃／ｈ以上で４５０～６００℃の温度範囲に加熱し、焼鈍温度：Ａｃ１変態点未満で１．０ｈ以上保持する焼鈍を施す高炭素熱延鋼板の製造方法。
［６］［１］～［４］のいずれかに記載の高炭素熱延鋼板の製造方法であって、前記成分組成を有する鋼を、熱間粗圧延後、仕上圧延終了温度：Ａｒ３変態点以上で仕上圧延を行い、その後、平均冷却速度：２０～１００℃／ｓｅｃで６５０～７５０℃まで冷却し、巻取温度：５００～７００℃で巻き取り、熱延鋼板とした後、該熱延鋼板を、平均加熱速度：１５℃／ｈ以上で４５０～６００℃の温度範囲に加熱し、Ａｃ１変態点以上Ａｃ３変態点以下で０．５ｈ以上保持し、次いで平均冷却速度：１～２０℃／ｈでＡｒ１変態点未満に冷却し、Ａｒ１変態点未満で２０ｈ以上保持する焼鈍を施す高炭素熱延鋼板の製造方法。 The present invention has been made based on the above findings, and has the following gist.
[1]% by mass, C: 0.20% or more and 0.50% or less, Si: 0.8% or less, Mn: 0.10% or more and 0.80% or less, P: 0.03% or less, S : 0.010% or less, sol. Al: 0.10% or less, N: 0.01% or less, Cr: 1.0% or less, B: 0.0005% or more and 0.005% or less, and one or two selected from Sb and Sn. The total content is 0.002% or more and 0.1% or less, and the balance has a composition of Fe and unavoidable impurities. The microstructure has an area ratio of 6 with respect to ferrite, cementite, and the whole microstructure. The ratio of the number of cementites having a circle-equivalent diameter of 0.1 μm or less to the total number of cementites is 20% or less, and the average cementite diameter is 2.5 μm or less. The area ratio of the cementite to the microstructure is 3.5% or more and 10.0% or less, the average concentration of the solid solution B in the region from the surface layer to the depth of 100 μm is 10 mass ppm or more, and the surface layer High carbon hot-rolled steel sheet having an average concentration of 70 mass ppm or less of N present as AlN in the region from the depth to 100 μm.
[2] The high carbon hot-rolled steel sheet according to [1], which has a tensile strength of 480 MPa or less and a total elongation of 33% or more.
[3] The high carbon hot-rolled steel sheet according to [1] or [2], wherein the ferrite has an average particle size of 4 to 25 μm.
[4] The high carbon according to any one of [1] to [3], which further contains, in mass%, one or two groups selected from the following Group A and Group B in addition to the above component composition. Hot rolled steel sheet.
Note Group A: Ti: 0.06% or less Group B: One or two or more selected from Nb, Mo, Ta, Ni, Cu, V, and W, respectively, 0.0005% or more 0.1 % Or less [5] The method for producing a high carbon hot-rolled steel sheet according to any one of [1] to [4], wherein the steel having the above-described composition is subjected to hot rough rolling, and then finish rolling finish temperature: Ar. After finish rolling at 3 or more transformation points, and then cooling at an average cooling rate of 20 to 100° C./sec to 650 to 750° C. and winding at a winding temperature of 500 to 700° C. to obtain a hot rolled steel sheet, A high-carbon hot-rolled steel sheet which is annealed by heating the hot-rolled steel sheet at an average heating rate of 15° C./h or more in a temperature range of 450 to 600° C. and holding it at an annealing temperature of less than Ac 1 transformation point for 1.0 hour or more. Manufacturing method.
[6] The method for producing a high-carbon hot-rolled steel sheet according to any one of [1] to [4], wherein the steel having the above-mentioned composition of ingredients is subjected to hot rough rolling, and then finish rolling finish temperature: Ar 3 transformation. Finish rolling is performed at a point or higher, then the average cooling rate: 20 to 100°C/sec is cooled to 650 to 750°C, and the coiling temperature is 500 to 700°C. The rolled steel sheet is heated to a temperature range of 450 to 600° C. at an average heating rate of 15° C./h or more, and is kept for 0.5 h or more at an Ac 1 transformation point or more and an Ac 3 transformation point or less, and then an average cooling rate: 1 to cooled to below Ar 1 transformation point at 20 ° C. / h, the method of producing a high-carbon hot-rolled steel sheet subjected to annealing for holding 20h or less than Ar 1 transformation point.
　本発明によれば、冷間加工性および焼入れ性（ズブ焼入れ性、浸炭焼入れ性）に優れた高炭素熱延鋼板を得られる。そして、本発明により製造した高炭素熱延鋼板を、素材鋼板として冷間加工性が必要とされるシートリクライナーやドアラッチ、および駆動系向けなどの自動車用部品に適用することにより、安定した品質が要求される自動車用部品の製造に大きく寄与でき、産業上格段の効果を奏する。 According to the present invention, it is possible to obtain a high carbon hot-rolled steel sheet excellent in cold workability and hardenability (dip hardenability, carburizing hardenability). Then, by applying the high-carbon hot-rolled steel sheet produced according to the present invention to a sheet recliner or a door latch that requires cold workability as a raw material steel sheet, and an automobile part such as for a drive system, stable quality is obtained. It can make a significant contribution to the manufacture of required automobile parts and has a marked industrial effect.
　以下に、本発明の高炭素熱延鋼板およびその製造方法について詳細に説明する。なお、本発明は以下の実施形態に限定されない。 The high carbon hot rolled steel sheet of the present invention and the manufacturing method thereof will be described in detail below. The present invention is not limited to the embodiments below.
　１）成分組成
　本発明の高炭素熱延鋼板の成分組成と、その限定理由について説明する。なお、以下の成分組成の含有量の単位である「％」は、特に断らない限り「質量％」を意味するものとする。 1) Component composition The component composition of the high carbon hot-rolled steel sheet of the present invention and the reason for limitation thereof will be described. In addition, "%" which is a unit of the content of the following component composition shall mean "mass %" unless there is particular notice.
　Ｃ：０．２０％以上０．５０％以下
　Ｃは、焼入れ後の強度を得るために重要な元素である。Ｃ量が０．２０％未満の場合、成形した後の熱処理によって所望の硬さが得られないため、Ｃ量は０．２０％以上にする必要がある。しかし、Ｃ量が０．５０％超えでは硬質化し、靭性や冷間加工性が劣化する。したがって、Ｃ量は０．２０％以上０．５０％以下とする。形状が複雑でプレス加工の難しい部品の冷間加工に用いる場合には、Ｃ量は０．４５％以下とすることが好ましく、０．４０％以下とすることがさらに好ましい。 C: 0.20% or more and 0.50% or less C is an important element for obtaining the strength after quenching. If the C content is less than 0.20%, the desired hardness cannot be obtained by the heat treatment after molding, so the C content needs to be 0.20% or more. However, if the amount of C exceeds 0.50%, it hardens and the toughness and cold workability deteriorate. Therefore, the C content is set to 0.20% or more and 0.50% or less. When used for cold working of a part having a complicated shape and difficult to press, the C content is preferably 0.45% or less, and more preferably 0.40% or less.
　Ｓｉ：０．８％以下
　Ｓｉは、固溶強化により強度を上昇させる元素である。Ｓｉ量の増加とともに硬質化し、冷間加工性が劣化するため、Ｓｉ量は０．８％以下とする。好ましくは０．６５％以下であり、さらに好ましくは０．５０％以下である。難成形部品用途において更なる冷間加工性が求められる場合には０．３０％以下とすることが好ましい。焼入れ後の焼き戻し工程で所定の軟化抵抗を確保するといった観点から、Ｓｉ量は、好ましくは０．１％以上とし、より好ましくは０．２％以上とする。 Si: 0.8% or less Si is an element that increases strength by solid solution strengthening. The amount of Si is 0.8% or less because it hardens as the amount of Si increases and the cold workability deteriorates. It is preferably 0.65% or less, more preferably 0.50% or less. When further cold workability is required in the use of difficult-to-form parts, it is preferably 0.30% or less. From the viewpoint of ensuring a predetermined softening resistance in the tempering process after quenching, the Si amount is preferably 0.1% or more, more preferably 0.2% or more.
　Ｍｎ：０．１０％以上０．８０％以下
　Ｍｎは、焼入れ性を向上させるとともに、固溶強化により強度を上昇させる元素である。Ｍｎ量が０．１０％未満になるとズブ焼入れ性および浸炭焼入れ性ともに低下し始めるため、Ｍｎ量は０．１０％以上とする。厚物材等で内部まで確実に焼入れる場合には、好ましくは０．２５％以上であり、さらに好ましくは０．３０％以上である。一方、Ｍｎ量が０．８０％を超えると、Ｍｎの偏析に起因したバンド組織が発達し、組織が不均一になり、かつ固溶強化により鋼が硬質化し冷間加工性が低下する。したがって、Ｍｎ量は０．８０％以下とする。成形性の求められる部品用の材料としては、所定の冷間加工性を必要とするため、Ｍｎ量は０．６５％以下とすることが好ましい。さらに好ましくは０．５５％以下である。 Mn: 0.10% or more and 0.80% or less Mn is an element that improves hardenability and increases strength by solid solution strengthening. If the Mn content is less than 0.10%, both the quench hardenability and the carburizing hardenability begin to deteriorate, so the Mn content is set to 0.10% or more. In the case of reliably quenching the inside of a thick material or the like, the content is preferably 0.25% or more, more preferably 0.30% or more. On the other hand, when the Mn content exceeds 0.80%, a band structure due to Mn segregation develops, the structure becomes nonuniform, and solid solution strengthens the steel to deteriorate the cold workability. Therefore, the amount of Mn is 0.80% or less. As a material for parts required to have moldability, a predetermined cold workability is required, so that the Mn content is preferably 0.65% or less. More preferably, it is 0.55% or less.
　Ｐ：０．０３％以下
　Ｐは、固溶強化により強度を上昇させる元素である。Ｐ量が０．０３％を超えて増加すると粒界脆化を招き、焼入れ後の靭性が劣化する。また、冷間加工性も低下させる。したがって、Ｐ量は０．０３％以下とする。優れた焼入れ後の靭性を得るには、Ｐ量は０．０２％以下が好ましい。Ｐは冷間加工性および焼入れ後の靭性を低下させるため、Ｐ量は少ないほど好ましい。しかしながら、過度にＰを低減すると精錬コストが増大するため、Ｐ量は０．００５％以上が好ましい。さらに好ましくは０．００７％以上である。 P: 0.03% or less P is an element that increases strength by solid solution strengthening. If the P content exceeds 0.03%, grain boundary embrittlement is caused, and the toughness after quenching deteriorates. Further, cold workability is also reduced. Therefore, the P content is 0.03% or less. In order to obtain excellent toughness after quenching, the P content is preferably 0.02% or less. Since P reduces the cold workability and the toughness after quenching, the smaller the amount of P, the more preferable. However, if the P content is excessively reduced, the refining cost increases, so the P content is preferably 0.005% or more. More preferably, it is 0.007% or more.
　Ｓ：０．０１０％以下
　Ｓは、硫化物を形成し、高炭素熱延鋼板の冷間加工性および焼入れ後の靭性を低下させるため、低減しなければならない元素である。Ｓ量が０．０１０％を超えると、高炭素熱延鋼板の冷間加工性および焼入れ後の靭性が著しく劣化する。したがって、Ｓ量は０．０１０％以下とする。優れた冷間加工性および焼入れ後の靭性を得るには、Ｓ量は０．００５％以下が好ましい。Ｓは、冷間加工性および焼入れ後の靭性を低下させるため、Ｓ量は少ないほど好ましい。しかしながら、過度にＳを低減すると精錬コストが増大するため、Ｓ量は０．０００５％以上が好ましい。 S: 0.010% or less S is an element that must be reduced because it forms a sulfide and reduces the cold workability and the toughness of the high carbon hot-rolled steel sheet after quenching. If the S content exceeds 0.010%, the cold workability and the toughness of the high carbon hot-rolled steel sheet after quenching are significantly deteriorated. Therefore, the S amount is 0.010% or less. In order to obtain excellent cold workability and toughness after quenching, the S content is preferably 0.005% or less. Since S lowers the cold workability and the toughness after quenching, it is preferable that the amount of S is smaller. However, if the S content is excessively reduced, the refining cost increases, so the S content is preferably 0.0005% or more.
　ｓｏｌ．Ａｌ：０．１０％以下
　ｓｏｌ．Ａｌ量が０．１０％を超えると、焼入れ処理の加熱時にＡｌＮが生成されてオーステナイト粒が微細化し過ぎる。これにより、冷却時にフェライト相の生成が促進され、ミクロ組織がフェライトとマルテンサイトとなり、焼入れ後の硬さが低下する。したがって、ｓｏｌ．Ａｌ量は、０．１０％以下とする。好ましくは０．０６％以下とする。なお、ｓｏｌ．Ａｌは、脱酸の効果を有しており、十分に脱酸するためには、０．００５％以上とすることが好ましい。 sol. Al: 0.10% or less sol. If the amount of Al exceeds 0.10%, AlN is generated during heating in the quenching treatment, and the austenite grains become too fine. As a result, the generation of the ferrite phase is promoted during cooling, the microstructure becomes ferrite and martensite, and the hardness after quenching decreases. Therefore, sol. The amount of Al is 0.10% or less. Preferably it is 0.06% or less. In addition, sol. Al has a deoxidizing effect and is preferably 0.005% or more for sufficient deoxidation.
　Ｎ：０．０１％以下
　Ｎ量が０．０１％を超えると、ＡｌＮの形成により焼入れ処理の加熱時にオーステナイト粒が微細化し過ぎ、冷却時にフェライト相の生成が促進され、焼入れ後の硬さが低下する。したがって、Ｎ量は、０．０１％以下とする。好ましくは０．００６５％以下である。さらに好ましくは、０．００５０％以下である。なお、Ｎは、ＡｌＮ、Ｃｒ系窒化物およびＢ窒化物を形成する。これにより、焼入れ処理の加熱時にオーステナイト粒の成長を適度に抑制して、焼入れ後の靭性を向上させる元素である。このため、Ｎ量は０．０００５％以上が好ましい。さらに好ましくは０．００１０％以上である。 N: 0.01% or less When the amount of N exceeds 0.01%, the austenite grains become too fine during heating in the quenching treatment due to the formation of AlN, the generation of a ferrite phase is promoted during cooling, and the hardness after quenching increases. descend. Therefore, the N content is 0.01% or less. It is preferably 0.0065% or less. More preferably, it is 0.0050% or less. Note that N forms AlN, a Cr-based nitride, and a B-nitride. This is an element that appropriately suppresses the growth of austenite grains during heating during the quenching treatment and improves the toughness after quenching. Therefore, the N content is preferably 0.0005% or more. More preferably, it is 0.0010% or more.
　Ｃｒ：１．０％以下
　本発明では、Ｃｒは、焼入れ性を高める重要な元素である。鋼中のＣｒ量が０％であると、特に浸炭焼入れにおいて表層でフェライトが発生しやすくなり、完全焼入れ組織が得られず、硬度低下が起こりやすい場合がある。このため、焼入れ性を重視する用途に用いる際には好ましくは０．０５％以上であり、さらに好ましくは０．１０％以上であり、より一層好ましくは０．２０％以上である。一方、Ｃｒ量が１．０％を超えると、焼入れ前の鋼板が硬質化して、冷間加工性が損なわれる。このため、Ｃｒ量は１．０％以下とする。なお、プレス成形の難しい高加工を必要とする部品を加工する際には、より一層優れた冷間加工性を必要とするため、Ｃｒ量は０．７％以下とすることが好ましく、０．５％以下とすることがさらに好ましい。 Cr: 1.0% or less In the present invention, Cr is an important element that enhances hardenability. When the Cr content in the steel is 0%, ferrite is likely to be generated in the surface layer particularly in the case of carburizing and quenching, a completely quenched structure cannot be obtained, and hardness is likely to be lowered. For this reason, when it is used in applications where hardenability is important, it is preferably at least 0.05%, more preferably at least 0.10%, and even more preferably at least 0.20%. On the other hand, if the Cr content exceeds 1.0%, the steel sheet before quenching becomes hard and the cold workability is impaired. Therefore, the Cr content is 1.0% or less. When processing a part that requires high workability, which is difficult to press-form, further excellent cold workability is required. Therefore, the Cr content is preferably 0.7% or less. It is more preferable to be 5% or less.
　Ｂ：０．０００５％以上０．００５％以下
　本発明では、Ｂは、焼入れ性を高める重要な元素である。Ｂ量が０．０００５％未満の場合、十分な効果が認められないため、Ｂ量は０．０００５％以上とする必要がある。好ましくは０．００１０％以上である。一方、Ｂ量が０．００５％超えの場合、仕上圧延後のオーステナイトの再結晶が遅延し、結果として熱延鋼板の集合組織が発達し、焼鈍後の異方性が大きくなり、絞り成形において耳が発生しやすくなる。このため、Ｂ量は０．００５％以下とする。好ましくは０．００４％以下である。 B: 0.0005% or more and 0.005% or less In the present invention, B is an important element that enhances hardenability. When the amount of B is less than 0.0005%, no sufficient effect is observed, so the amount of B must be 0.0005% or more. It is preferably 0.0010% or more. On the other hand, when the B content is more than 0.005%, recrystallization of austenite after finish rolling is delayed, resulting in the development of texture of the hot rolled steel sheet, the anisotropy after annealing becomes large, and in draw forming. Ears are more likely to occur. Therefore, the B content is 0.005% or less. It is preferably 0.004% or less.
　ＳｎおよびＳｂから選んだ１種または２種の合計：０．００２％以上０．１％以下
　Ｓｂ、Ｓｎは、鋼板表層からの浸窒抑制に有効な元素である。これら元素の１種以上の合計が０．００２％未満の場合、十分な効果が認められないため、これら元素の１種以上の合計は０．００２％以上とする。さらに好ましくは０．００５％以上である。一方、これらの元素の１種以上の合計が０．１％を超えて含有しても、浸窒防止効果は飽和する。また、これらの元素は、粒界に偏析する傾向があるため、合計で０．１％超えとすると、含有量が高くなりすぎ、粒界脆化を引き起こす可能性がある。したがって、ＳｂおよびＳｎのうちから選んだ１種または２種の合計の含有量は、０．１％以下とする。好ましくは０．０３％以下であり、さらに好ましくは０．０２％以下である。 Total of one or two selected from Sn and Sb: 0.002% or more and 0.1% or less Sb and Sn are elements effective for suppressing nitriding from the steel sheet surface layer. If the total of one or more of these elements is less than 0.002%, a sufficient effect is not observed, so the total of one or more of these elements is set to 0.002% or more. More preferably, it is 0.005% or more. On the other hand, even if the total content of one or more of these elements exceeds 0.1%, the effect of preventing nitrification is saturated. Further, since these elements tend to segregate at the grain boundaries, if the total content exceeds 0.1%, the content becomes too high, which may cause grain boundary embrittlement. Therefore, the total content of one or two selected from Sb and Sn is 0.1% or less. It is preferably 0.03% or less, more preferably 0.02% or less.
　本発明では、ＳｂおよびＳｎのうちから選んだ１種または２種を合計で０．００２％以上０．１％以下とすることで、窒素雰囲気で焼鈍した場合でも鋼板表層からの浸窒を抑制し、鋼板表層における窒素濃度の増加を抑制する。このように、本発明によれば、鋼板表層からの浸窒を抑制できるため、窒素雰囲気で焼鈍した場合であっても、焼鈍後の鋼板表層から深さ１００μｍまでの領域における固溶Ｂ量を適切に確保することができ、かつ鋼板表層から深さ１００μｍまでの領域におけるＡｌ窒化物（ＡｌＮ）の生成を抑えることで焼入れ前加熱時のオーステナイト粒が成長できる。その結果、冷却時にフェライトおよびパーライトの生成を遅らせることができるため、これにより高い焼入れ性を得ることができる。 In the present invention, by controlling the total of one or two selected from Sb and Sn to be 0.002% or more and 0.1% or less, it is possible to suppress nitriding from the surface layer of the steel sheet even when annealed in a nitrogen atmosphere. However, the increase in nitrogen concentration in the surface layer of the steel sheet is suppressed. As described above, according to the present invention, since the nitriding from the steel sheet surface layer can be suppressed, the amount of solid solution B in the region from the steel sheet surface layer after annealing to a depth of 100 μm can be suppressed even when annealed in a nitrogen atmosphere. It can be properly secured, and by suppressing the generation of Al nitride (AlN) in the region from the surface of the steel sheet to a depth of 100 μm, austenite grains at the time of heating before quenching can grow. As a result, the production of ferrite and pearlite can be delayed during cooling, and thus high hardenability can be obtained.
　本発明において、上記以外の残部は、Ｆｅおよび不可避的不純物である。 In the present invention, the balance other than the above is Fe and inevitable impurities.
　以上の必須含有元素で、本発明の高炭素熱延鋼板は目的とする特性が得られる。なお、本発明の高炭素熱延鋼板は、例えば焼入れ性をさらに向上させることを目的として、必要に応じて下記の元素を含有することができる。 With the above essential elements, the high-carbon hot-rolled steel sheet of the present invention can obtain the desired characteristics. The high-carbon hot-rolled steel sheet of the present invention may contain the following elements, if necessary, for the purpose of further improving hardenability.
　Ｔｉ：０．０６％以下
　Ｔｉは、焼入れ性を高めるために有効な元素である。Ｂの含有のみでは焼入れ性が不十分な場合に、Ｔｉを含有することで、焼入れ性を向上させることができる。Ｔｉ量が０．００５％未満では、その効果が認められないため、Ｔｉを含有する場合、Ｔｉ量は０．００５％以上とすることが好ましい。さらに好ましくは０．００７％以上である。一方、Ｔｉ量が０．０６％を超えて含有すると、焼入れ前の鋼板が硬質化して冷間加工性が損なわれるため、Ｔｉを含有する場合、Ｔｉ量は０．０６％以下とする。好ましくは０．０４％以下である。 Ti: 0.06% or less Ti is an element effective for improving hardenability. When the hardenability is insufficient only by containing B, the hardenability can be improved by containing Ti. If the Ti content is less than 0.005%, the effect is not recognized. Therefore, when Ti is contained, the Ti content is preferably 0.005% or more. More preferably, it is 0.007% or more. On the other hand, if the Ti content exceeds 0.06%, the steel sheet before quenching becomes hard and the cold workability is impaired. Therefore, when Ti is contained, the Ti content is 0.06% or less. It is preferably 0.04% or less.
　さらに、本発明の機械特性および焼入れ性を安定化させるためにＮｂ、Ｍｏ、Ｔａ、Ｎｉ、Ｃｕ、Ｖ、Ｗのうちから選んだ１種または２種以上を、それぞれ所要量、添加してもよい。 Furthermore, in order to stabilize the mechanical properties and hardenability of the present invention, one or more selected from Nb, Mo, Ta, Ni, Cu, V, and W may be added in the required amounts, respectively. Good.
　Ｎｂ：０．０００５％以上０．１％以下
　Ｎｂは、炭窒化物を形成し、焼入れ前加熱時の結晶粒の異常粒成長の防止や靱性改善、焼戻し軟化抵抗改善に有効な元素である。０．０００５％未満では添加効果は十分に発現しないため、Ｎｂを含有する場合には下限を０．０００５％とすることが好ましい。さらに好ましくは０．００１０％以上とする。Ｎｂは０．１％を超えると添加効果が飽和するだけでなく、Ｎｂ炭化物により母材の引張強度の増加に伴い伸びを低下させることになるため、Ｎｂを含有する場合には上限を０．１％とすることが好ましい。さらに好ましくは０．０５％以下であり、より一層好ましくは０．０３％未満である。 Nb: 0.0005% or more and 0.1% or less Nb is an element that forms carbonitrides and is effective in preventing abnormal grain growth of crystal grains during heating before quenching, improving toughness, and improving temper softening resistance. If it is less than 0.0005%, the effect of addition is not sufficiently exhibited, so when Nb is contained, the lower limit is preferably made 0.0005%. More preferably, it is 0.0010% or more. If Nb exceeds 0.1%, not only the effect of addition is saturated, but also Nb carbides reduce the elongation as the tensile strength of the base material increases. Therefore, when Nb is contained, the upper limit is set to 0. It is preferably set to 1%. It is more preferably 0.05% or less, and even more preferably less than 0.03%.
　Ｍｏ：０．０００５％以上０．１％以下
　Ｍｏは、焼入れ性の向上と、焼戻し軟化抵抗性の向上に有効な元素である。０．０００５％未満では添加効果が小さいので、Ｍｏを含有する場合には下限を０．０００５％とすることが好ましい。さらに好ましくは０．００１０％以上とする。Ｍｏは０．１％を超えると添加効果は飽和し、コストも増加するため、Ｍｏを含有する場合には上限を０．１％とすることが好ましい。さらに好ましくは、０．０５％以下であり、より一層好ましくは０．０３％未満である。 Mo: 0.0005% or more and 0.1% or less Mo is an element effective for improving hardenability and temper softening resistance. If less than 0.0005%, the effect of addition is small. Therefore, when Mo is contained, the lower limit is preferably 0.0005%. More preferably, it is 0.0010% or more. When Mo exceeds 0.1%, the effect of addition is saturated and the cost also increases. Therefore, when Mo is contained, the upper limit is preferably 0.1%. It is more preferably 0.05% or less, and even more preferably less than 0.03%.
　Ｔａ：０．０００５％以上０．１％以下
　Ｔａは、Ｎｂと同様に炭窒化物を形成し、焼入れ前加熱時の結晶粒の異常粒成長防止や結晶粒の粗大化防止、焼戻し軟化抵抗改善に有効な元素である。０．０００５％未満では添加効果が小さいので、Ｔａを含有する場合には下限を０．０００５％とすることが好ましい。さらに好ましくは０．００１０％以上とする。Ｔａは０．１％を超えると添加効果が飽和したり、過剰な炭化物形成による焼入れ硬度を低下させたり、またコスト増となるため、Ｔａを含有する場合には上限を０．１％とすることが好ましい。さらに好ましくは、０．０５％以下であり、より一層好ましくは０．０３％未満である。 Ta: 0.0005% or more and 0.1% or less Ta forms carbonitrides like Nb, prevents abnormal grain growth of crystal grains during heating before quenching, prevents crystal grain coarsening, and improves temper softening resistance. Is an effective element. If less than 0.0005%, the effect of addition is small. Therefore, when Ta is contained, the lower limit is preferably 0.0005%. More preferably, it is 0.0010% or more. If Ta exceeds 0.1%, the effect of addition is saturated, quenching hardness is reduced due to excessive carbide formation, and the cost is increased. Therefore, when Ta is contained, the upper limit is 0.1%. It is preferable. It is more preferably 0.05% or less, and even more preferably less than 0.03%.
　Ｎｉ：０．０００５％以上０．１％以下
　Ｎｉは靱性の向上や焼入れ性の向上に効果の高い元素である。０．０００５％未満では添加効果がないため、Ｎｉを含有する場合には下限を０．０００５％とすることが好ましい。さらに好ましくは０．００１０％以上とする。Ｎｉは０．１％超では、添加効果が飽和する上にコスト増加も招くため、Ｎｉを含有する場合には上限を０．１％とすることが好ましい。さらに好ましくは、０．０５％以下である。 Ni: 0.0005% or more and 0.1% or less Ni is an element highly effective in improving toughness and hardenability. If less than 0.0005%, there is no effect of addition, so when Ni is contained, the lower limit is preferably 0.0005%. More preferably, it is 0.0010% or more. If Ni exceeds 0.1%, the effect of addition is saturated and the cost also increases. Therefore, when Ni is contained, the upper limit is preferably made 0.1%. More preferably, it is 0.05% or less.
　Ｃｕ：０．０００５％以上０．１％以下
　Ｃｕは、焼入れ性の確保に有効な元素である。０．０００５％未満では添加効果が十分に確認されないため、Ｃｕを含有する場合には下限を０．０００５％とすることが好ましい。さらに好ましくは０．００１０％以上とする。Ｃｕは０．１％超では、熱延時の疵が発生しやすくなり歩留りを落とすなど製造性を劣化させるので、Ｃｕを含有する場合には上限を０．１％とすることが好ましい。さらに好ましくは、０．０５％以下である。 Cu: 0.0005% or more and 0.1% or less Cu is an element effective for ensuring hardenability. If less than 0.0005%, the effect of addition is not sufficiently confirmed. Therefore, when Cu is contained, the lower limit is preferably 0.0005%. More preferably, it is 0.0010% or more. If Cu is more than 0.1%, defects during hot rolling tend to occur and the productivity is deteriorated such as a decrease in yield. Therefore, when Cu is contained, the upper limit is preferably 0.1%. More preferably, it is 0.05% or less.
　Ｖ：０．０００５％以上０．１％以下
　Ｖは、ＮｂやＴａと同様に、炭窒化物を形成し、焼入れ前加熱時の結晶粒の異常粒成長防止および靱性改善、焼戻し軟化抵抗改善に有効な元素である。０．０００５％未満では添加効果は十分に発現しないため、Ｖを含有する場合には下限を０．０００５％とすることが好ましい。さらに好ましくは０．００１０％以上とする。Ｖは０．１％を超えると添加効果が飽和するだけでなく、Ｎｂ炭化物により母材の引張強度の増加に伴い伸びを低下させることになるため、Ｖを含有する場合には上限を０．１％とすることが好ましい。さらに好ましくは、０．０５％以下であり、より一層好ましくは０．０３％未満である。 V: 0.0005% or more and 0.1% or less V, like Nb and Ta, forms carbonitrides to prevent abnormal grain growth of crystal grains during heating before quenching, improve toughness, and improve temper softening resistance. It is an effective element. If it is less than 0.0005%, the effect of addition is not sufficiently exhibited, so when V is contained, the lower limit is preferably made 0.0005%. More preferably, it is 0.0010% or more. If V exceeds 0.1%, not only the effect of addition is saturated, but also the elongation decreases as the tensile strength of the base material increases due to Nb carbide. Therefore, when V is contained, the upper limit is set to 0. It is preferably set to 1%. It is more preferably 0.05% or less, and even more preferably less than 0.03%.
　Ｗ：０．０００５％以上０．１％以下
　Ｗは、Ｎｂ、Ｖと同様に、炭窒化物を形成し、焼入れ前加熱時のオーステナイト結晶粒の異常粒成長防止や焼き戻し軟化抵抗改善に有効な元素である。０．０００５％未満では添加効果が小さいので、Ｗを含有する場合には下限を０．０００５％とすることが好ましい。さらに好ましくは０．００１０％以上とする。Ｗは０．１％を超えると添加効果が飽和したり、過剰な炭化物形成による焼入れ硬度を低下させたり、またコスト増となるため、Ｗを含有する場合には上限を０．１％とすることが好ましい。さらに好ましくは、０．０５％以下であり、より一層好ましくは０．０３％未満である。 W: 0.0005% or more and 0.1% or less W, like Nb and V, forms carbonitrides and is effective in preventing abnormal grain growth of austenite crystal grains during heating before quenching and improving temper softening resistance. Is an element. If it is less than 0.0005%, the effect of addition is small, so when W is contained, the lower limit is preferably 0.0005%. More preferably, it is 0.0010% or more. If W exceeds 0.1%, the effect of addition is saturated, the quenching hardness is reduced due to excessive carbide formation, and the cost increases, so the upper limit is made 0.1% when W is contained. It is preferable. It is more preferably 0.05% or less, and even more preferably less than 0.03%.
　なお、本発明では、Ｎｂ、Ｍｏ、Ｔａ、Ｎｉ、Ｃｕ、Ｖ、Ｗのうちから選んだ２種以上を含有する場合には、その合計量を０．００１％以上０．１％以下とすることが好ましい。 In the present invention, when two or more kinds selected from Nb, Mo, Ta, Ni, Cu, V and W are contained, the total amount is made 0.001% or more and 0.1% or less. It is preferable.
　２）ミクロ組織
　本発明の高炭素熱延鋼板のミクロ組織の限定理由について説明する。 2) Microstructure The reason for limiting the microstructure of the high carbon hot-rolled steel sheet of the present invention will be described.
　本発明では、ミクロ組織は、フェライトおよびセメンタイトを有し、該セメンタイトは、円相当直径が０．１μｍ以下のセメンタイト数が全セメンタイト数に対して２０％以下、平均セメンタイト径は２．５μｍ以下、全ミクロ組織に対する上記セメンタイトの占める割合が面積率で３．５％以上１０．０％以下であり、表層から深さ１００μｍまでの領域における固溶Ｂ量の平均濃度が１０質量ｐｐｍ以上であり、表層から深さ１００μｍまでの領域におけるＡｌＮとして存在するＮ量の平均濃度が７０質量ｐｐｍ以下である。
また、本発明において、フェライトの平均粒径は４～２５μｍであることが好ましい。より好ましくは５μｍ以上である。 In the present invention, the microstructure has ferrite and cementite, and the cementite has a circle equivalent diameter of 0.1 μm or less with respect to the total cementite number of 20% or less, and an average cementite diameter of 2.5 μm or less, The area ratio of the cementite to the total microstructure is 3.5% or more and 10.0% or less, and the average concentration of the solid solution B in the region from the surface layer to the depth of 100 μm is 10 mass ppm or more, The average concentration of N present as AlN in the region from the surface layer to a depth of 100 μm is 70 mass ppm or less.
Further, in the present invention, it is preferable that the average particle diameter of the ferrite is 4 to 25 μm. More preferably, it is 5 μm or more.
　２－１）フェライトおよびセメンタイト
　本発明の高炭素熱延鋼板のミクロ組織は、フェライトおよびセメンタイトを有する。なお、本発明において、フェライトは面積率で９０％以上が好ましい。フェライト面積率が９０％未満となると成形性が悪くなり、加工度の高い部品で冷間加工が難しくなる場合がある。そのため、フェライト面積率は９０％以上が好ましい。さらに好ましくは９２％以上とする。 2-1) Ferrite and Cementite The microstructure of the high carbon hot-rolled steel sheet of the present invention has ferrite and cementite. In the present invention, the area ratio of ferrite is preferably 90% or more. If the ferrite area ratio is less than 90%, the formability is deteriorated, and cold working may be difficult for parts with high workability. Therefore, the ferrite area ratio is preferably 90% or more. More preferably, it is 92% or more.
　なお、本発明の高炭素熱延鋼板のミクロ組織は、上記したフェライトとセメンタイト以外に、パーライトが生成されてもよい。全ミクロ組織に対するパーライトの面積率が６．５％以下であれば、本発明の効果を損ねるものではないため、含有しても構わない。 Note that, in the microstructure of the high carbon hot-rolled steel sheet of the present invention, pearlite may be generated in addition to the above-mentioned ferrite and cementite. If the area ratio of pearlite with respect to the entire microstructure is 6.5% or less, the effect of the present invention is not impaired, and thus it may be included.
　２－２）全セメンタイト数に対する円相当直径０．１μｍ以下のセメンタイト数の割合：２０％以下
　円相当直径が０．１μｍ以下のセメンタイトが多いと分散強化により硬質化し、伸びが低下する。冷間加工性を得る観点より、本発明では、円相当直径が０．１μｍ以下のセメンタイト数を、全セメンタイト数に対して２０％以下とする。その結果、さらに、引張強度で４８０ＭＰａ以下、全伸び（Ｅｌ）が３３％以上を達成することができる。
難成形部品に用いる場合には高い冷間加工性が必要であり、この場合には、円相当直径が０．１μｍ以下のセメンタイト数が、全セメンタイト数に対して１０％以下であることが好ましい。円相当直径が０．１μｍ以下のセメンタイト数を、全セメンタイト数に対して１０％以下とすることで、引張強度で４４０ＭＰａ以下、全伸び（Ｅｌ）が３６％以上を達成することができる。なお、円相当直径が０．１μｍ以下のセメンタイトの割合を定義した理由は、０．１μｍ以下のセメンタイトでは分散強化能を生じ、その大きさのセメンタイトが増えると冷間加工性に支障をきたすためである。
焼鈍中におけるフェライト粒の異常粒成長抑制の観点から、円相当直径が０．１μｍ以下のセメンタイト数を、全セメンタイト数に対して３％以上とすることが好ましい。 2-2) Ratio of the number of cementites having a circle-equivalent diameter of 0.1 μm or less to the total number of cementite: 20% or less If there is a large amount of cementite having a circle-equivalent diameter of 0.1 μm or less, it becomes hardened due to dispersion strengthening and elongation is reduced. From the viewpoint of obtaining cold workability, in the present invention, the number of cementites having a circle equivalent diameter of 0.1 μm or less is 20% or less with respect to the total number of cementites. As a result, it is possible to further achieve a tensile strength of 480 MPa or less and a total elongation (El) of 33% or more.
High cold workability is required for use in difficult-to-form parts. In this case, the number of cementites having a circle equivalent diameter of 0.1 μm or less is preferably 10% or less of the total number of cementites. .. By setting the number of cementites having a circle equivalent diameter of 0.1 μm or less to 10% or less with respect to the total number of cementites, it is possible to achieve a tensile strength of 440 MPa or less and a total elongation (El) of 36% or more. The reason for defining the proportion of cementite having a circle-equivalent diameter of 0.1 μm or less is that cementite having a diameter of 0.1 μm or less produces dispersion strengthening ability, and if the size of cementite increases, cold workability is impaired. Is.
From the viewpoint of suppressing abnormal grain growth of ferrite grains during annealing, the number of cementites having a circle equivalent diameter of 0.1 μm or less is preferably 3% or more with respect to the total number of cementites.
　なお、焼入れ前に存在するセメンタイト径は、円相当直径で０．０７～３．０μｍ程度である。そのため、析出強化にそれほど影響しないサイズである、焼入れ前の円相当直径が０．１μｍ超のセメンタイトの分散状態については、特に本発明では規定しない。 Note that the cementite diameter existing before quenching is about 0.07 to 3.0 μm in equivalent circle diameter. Therefore, the dispersed state of cementite having a circle equivalent diameter of more than 0.1 μm before quenching, which is a size that does not significantly affect precipitation strengthening, is not particularly specified in the present invention.
　２－３）平均セメンタイト径：２．５μｍ以下
　焼入れ時にはセメンタイトを全て溶かして、所定のフェライト中の固溶Ｃ量を確保する必要がある。平均セメンタイト径が２．５μｍを超えるとオーステナイト域での保持中においてセメンタイトが完全に溶解できないため、平均セメンタイト径は２．５μｍ以下とする。より好ましくは２．０μｍ以下である。なお、セメンタイトが微細すぎるとセメンタイトの析出強化により冷間加工性が低下するため、平均セメンタイト径は０．１μｍ以上が好ましい。さらに好ましくは０．１５μｍ以上とする。
　なお、本発明において「セメンタイト径」とはセメンタイトの円相当直径を指し、セメンタイトの円相当直径は、セメンタイトの長径と短径を測定して円相当直径に換算した値とする。また「平均セメンタイト径」とは、円相当直径に換算した全てのセメンタイトの円相当直径の合計を、セメンタイト総数で除して求めた値を指す。 2-3) Average cementite diameter: 2.5 μm or less At the time of quenching, it is necessary to melt all the cementite to secure a predetermined amount of solid solution C in ferrite. If the average cementite diameter exceeds 2.5 μm, the cementite cannot be completely dissolved during holding in the austenite region, so the average cementite diameter is set to 2.5 μm or less. It is more preferably 2.0 μm or less. If the cementite is too fine, the precipitation strengthening of the cementite deteriorates the cold workability. Therefore, the average cementite diameter is preferably 0.1 μm or more. More preferably, it is 0.15 μm or more.
In the present invention, the “cementite diameter” refers to a circle-equivalent diameter of cementite, and the circle-equivalent diameter of cementite is a value obtained by measuring the major axis and the minor axis of cementite and converting them to the circle-equivalent diameters. Further, the “average cementite diameter” refers to a value obtained by dividing the sum of the equivalent circle diameters of all the cementites converted into equivalent circle diameters by the total number of cementites.
　２－４）全ミクロ組織に対するセメンタイトの占める割合（面積率）：３．５％以上１０．０％以下
　全ミクロ組織に対するセメンタイトの割合が１０．０％超えになると、それに伴い、析出強化に寄与する０．１μｍ以下のセメンタイト数も増加し、鋼が硬質化するため、１０．０％以下とする。好ましくは９．５％以下である。一方、上記割合が３．５％未満になると実質的なＣ含有量が０．２０％に達せず、熱処理後に所定の硬さが得られないため、３．５％以上とする。さらに好ましくは４．０％以上とする。 2-4) Proportion (area ratio) of cementite to all microstructures: 3.5% or more and 10.0% or less When the ratio of cementite to all microstructures exceeds 10.0%, it contributes to precipitation strengthening. The number of cementite particles having a particle size of 0.1 μm or less increases, and the steel hardens, so the content is made 10.0% or less. It is preferably 9.5% or less. On the other hand, if the above ratio is less than 3.5%, the substantial C content does not reach 0.20%, and the predetermined hardness cannot be obtained after heat treatment, so the content is made 3.5% or more. More preferably, it is 4.0% or more.
　２－５）フェライトの平均粒径：４～２５μｍ（好適条件）
　フェライトの平均粒径は、４μｍ未満では冷間加工前の強度が増加し、プレス成形性が劣化する恐れがあるため、４μｍ以上が好ましい。一方、フェライトの平均粒径は２５μｍを超えると、母材強度が低下する恐れがある。また、目的とする製品形状に成型加工後、焼入れせずに使用する領域では、ある程度母材の強度が必要である。そのため、フェライト平均粒径は、２５μｍ以下とすることが好ましい。さらに好ましくは５μｍ以上、より一層好ましくは６μｍ以上である。さらに好ましくは２０μｍ以下である。より一層好ましくは１８μｍ以下である。 2-5) Average particle size of ferrite: 4 to 25 μm (suitable condition)
If the average particle size of ferrite is less than 4 μm, the strength before cold working increases and the press formability may deteriorate, so 4 μm or more is preferable. On the other hand, if the average grain size of ferrite exceeds 25 μm, the strength of the base material may decrease. Further, the strength of the base material is required to some extent in a region where it is used without being hardened after being molded into a desired product shape. Therefore, the average ferrite grain size is preferably 25 μm or less. It is more preferably at least 5 μm, and even more preferably at least 6 μm. More preferably, it is 20 μm or less. More preferably, it is 18 μm or less.
　なお、本発明では、上述のセメンタイトの円相当直径、平均セメンタイト径、全ミクロ組織に対するセメンタイトの占める割合、フェライトの面積率、フェライトの平均粒径等は、後述する実施例に記載の方法で測定することができる。 In the present invention, the equivalent circle diameter of cementite, the average cementite diameter, the ratio of cementite to the total microstructure, the area ratio of ferrite, the average particle diameter of ferrite, etc. are measured by the method described in Examples described later. can do.
　２－６）表層から深さ１００μｍまでの領域における固溶Ｂ量の平均濃度：１０質量ｐｐｍ以上
　本発明の高炭素熱延鋼板においては、鋼板を焼入れした際に表層部に生成しやすいパーライト、ソルバイトといわれるような焼入れ組織を防止するために、鋼板表層から板厚方向へ１００μｍ位置までの領域（部位）（表層１００μｍ部）のＢ量が、窒化や酸化していない固溶Ｂとして平均濃度で１０質量ｐｐｍ以上存在する必要がある。焼入れ処理を行って使用する耐摩耗性が必要とされる自動車部品では表面硬度が要求される。所定の表面硬度を得るためには焼入れ後表層１００μｍ部において完全焼入れ組織を得る必要がある。好ましくは、上記固溶Ｂ量の平均濃度は１２質量ｐｐｍ以上である。さらに好ましくは１５質量ｐｐｍ以上である。なお、固溶Ｂが高すぎると熱延組織の集合組織の発達の妨げになるため、４０質量ｐｐｍ以下とする。さらに好ましくは３５質量ｐｐｍ以下とする。 2-6) Average concentration of solid solution B in the region from the surface layer to a depth of 100 μm: 10 mass ppm or more In the high carbon hot-rolled steel sheet of the present invention, pearlite which is easily generated in the surface layer portion when the steel sheet is quenched, In order to prevent a quenching structure called sorbite, the amount of B in the region (portion) from the surface of the steel plate to 100 μm in the thickness direction (100 μm surface) is the average concentration as solid solution B that has not been nitrided or oxidized. It is necessary to exist at 10 mass ppm or more. Surface hardness is required for automobile parts that require wear resistance after being subjected to quenching treatment. In order to obtain a predetermined surface hardness, it is necessary to obtain a completely quenched structure in the surface layer of 100 μm after quenching. Preferably, the average concentration of the solute B is 12 mass ppm or more. More preferably, it is 15 mass ppm or more. If the solid solution B is too high, the development of the texture of the hot rolled structure is hindered, so the content is set to 40 mass ppm or less. More preferably, it is 35 mass ppm or less.
　２－７）表層から深さ１００μｍまでの領域におけるＡｌＮとして存在するＮ量の平均濃度：７０質量ｐｐｍ以下
　鋼板表層から板厚方向へ１００μｍ位置までの領域におけるＡｌＮとして存在するＮ量の平均濃度を７０質量ｐｐｍ以下とすることで、焼入れ前加熱におけるオーステナイト域で結晶粒の成長を促進する。これにより、冷却段階でパーライト、ソルバイトといわれる組織が得られにくくなり、焼き入れ不足が起こらず、所定の表面硬度が得られる。表層から深さ１００μｍまでの領域におけるＡｌＮとして存在するＮ量の平均濃度は５０質量ｐｐｍ以下とすることが好ましい。
なお、オーステナイト域での加熱において異常粒成長を抑制する観点から、上記Ｎ量の平均濃度は、１０質量ｐｐｍ以上とすることが好ましく、２０質量ｐｐｍ以上とすることがさらに好ましい。 2-7) Average concentration of N amount existing as AlN in the region from the surface layer to the depth of 100 μm: 70 mass ppm or less The average concentration of N amount present as AlN in the region from the steel plate surface layer to the 100 μm position in the plate thickness direction When the content is 70 mass ppm or less, the growth of crystal grains is promoted in the austenite region in the heating before quenching. This makes it difficult to obtain a structure called pearlite or sorbite in the cooling stage, does not cause insufficient quenching, and has a predetermined surface hardness. The average concentration of the amount of N existing as AlN in the region from the surface layer to the depth of 100 μm is preferably 50 mass ppm or less.
In addition, from the viewpoint of suppressing abnormal grain growth during heating in the austenite region, the average concentration of the N content is preferably 10 mass ppm or more, and more preferably 20 mass ppm or more.
　本発明では、鋼板表層部における固溶Ｂ量およびＡｌＮとして存在するＮ量は、加熱条件、巻取条件、焼鈍条件の各工程での製造条件が密接に関係し、これらの一連の製造条件を最適化することが必要であることが判明した。なお、各工程で固溶Ｂ量およびＡｌＮとしてのＮ量を得るために必要な理由は後述する。 In the present invention, the amount of solid solution B in the steel sheet surface layer and the amount of N present as AlN are closely related to the manufacturing conditions in each step such as heating conditions, winding conditions, and annealing conditions. It turned out to be necessary to optimize. The reason necessary to obtain the amount of solid solution B and the amount of N as AlN in each step will be described later.
　３）機械特性
　本発明の高炭素熱延鋼板は、ギア、トランスミッション、シートリクライナーなどの自動車用部品を冷間プレスで成形するため、優れた冷間加工性が必要である。また、焼入れ処理により硬さを大きくして、耐磨耗性を付与する必要がある。そのため、本発明の高炭素熱延鋼板は、鋼板の引張強度を低減して引張強度（ＴＳ）を４８０ＭＰａ以下とし、かつ伸びを高めて全伸び（Ｅｌ）を３３％以上とすることで、優れた冷間加工性を有するとともに、優れた焼入れ性（ズブ焼入れ性、浸炭焼入れ性）を両立させることができる。さらに好ましくは、ＴＳを４６０ＭＰａ以下とし、Ｅｌを３５％以上とする。 3) Mechanical Properties The high-carbon hot-rolled steel sheet of the present invention is required to have excellent cold workability because it is used to form automobile parts such as gears, transmissions, and seat recliners by cold pressing. Further, it is necessary to increase hardness by quenching treatment to impart wear resistance. Therefore, the high carbon hot-rolled steel sheet of the present invention is excellent by reducing the tensile strength of the steel sheet to a tensile strength (TS) of 480 MPa or less and increasing the elongation to a total elongation (El) of 33% or more. It has both cold workability and excellent hardenability (dip hardenability, carburizing hardenability). More preferably, TS is 460 MPa or less and El is 35% or more.
　また、冷間プレス性を必要とする難成形部品を成形することを想定して、さらに鋼板の引張強度を低減してＴＳを４４０ＭＰａ以下とし、かつ全伸びを高めてＥｌを３６％以上とすることで優れた冷間加工性を有するとともに、優れた焼入れ性（ズブ焼入れ性、浸炭焼入れ性）を両立することができる。さらに好ましくはＴＳを４１０ＭＰａ以下とし、Ｅｌを３８％以上とする。 In addition, assuming that a difficult-to-form part requiring cold pressability is formed, the tensile strength of the steel sheet is further reduced to TS of 440 MPa or less, and the total elongation is increased to El of 36% or more. This makes it possible to have both excellent cold workability and excellent hardenability (dip hardenability, carburizing hardenability). More preferably, TS is 410 MPa or less and El is 38% or more.
　なお、上述の引張強度（ＴＳ）、全伸び（Ｅｌ）は、後述する実施例に記載の方法で測定することができる。 The above-mentioned tensile strength (TS) and total elongation (El) can be measured by the methods described in Examples below.
　４）製造方法
　本発明の高炭素熱延鋼板は、上記のような成分組成を有する鋼を素材とし、この素材（鋼素材）を熱間粗圧延後、仕上圧延終了温度：Ａｒ３変態点以上で仕上圧延を行い、その後、平均冷却速度：２０～１００℃／ｓｅｃで６５０～７５０℃まで冷却し、巻取温度：５００～７００℃で巻き取り、常温まで冷却して熱延鋼板とした後、熱延鋼板を、平均加熱速度：１５℃／ｈ以上で４５０～６００℃の温度範囲に加熱し、焼鈍温度：Ａｃ１変態点未満で１．０ｈ以上保持する焼鈍を施すことにより製造される。 4) Manufacturing method The high-carbon hot-rolled steel sheet of the present invention is made of steel having the above-described composition, and after this material (steel material) is hot-roughly rolled, finish rolling end temperature: Ar 3 transformation point or higher. After finish rolling, the average cooling rate: 20-100°C/sec, cooling to 650-750°C, winding temperature: 500-700°C, cooling to room temperature to obtain a hot rolled steel sheet Manufactured by heating a hot rolled steel sheet at an average heating rate of 15° C./h or more in a temperature range of 450 to 600° C. and annealing at a temperature of less than Ac 1 transformation point for 1.0 h or more. ..
　または、上記のような成分組成を有する鋼を素材とし、この素材（鋼素材）を熱間粗圧延後、仕上圧延終了温度：Ａｒ３変態点以上で仕上圧延を行い、その後、平均冷却速度：２０～１００℃／ｓｅｃで６５０～７５０℃まで冷却し、巻取温度：５００～７００℃で巻き取り、常温まで冷却して熱延鋼板とした後、熱延鋼板を、平均加熱速度：１５℃／ｈ以上で４５０～６００℃の温度範囲に加熱し、Ａｃ１変態点以上Ａｃ３変態点以下で０．５ｈ以上保持し、次いで平均冷却速度：１～２０℃／ｈでＡｒ１変態点未満に冷却し、Ａｒ１変態点未満で２０ｈ以上保持する２段焼鈍を施すことにより製造される。 Alternatively, using a steel having the above-described composition as a raw material, this raw material (steel raw material) is subjected to hot rough rolling, then finish rolling at a finish rolling end temperature: Ar 3 transformation point or higher, and then an average cooling rate: After cooling to 650 to 750° C. at 20 to 100° C./sec, winding at a winding temperature of 500 to 700° C. and cooling to room temperature to form a hot rolled steel sheet, the hot rolled steel sheet has an average heating rate of 15° C. /H or more and heating to a temperature range of 450 to 600° C., maintaining at Ac 1 transformation point or more and Ac 3 transformation point or less for 0.5 h or more, and then average cooling rate: 1 to 20° C./h and less than Ar 1 transformation point And is subjected to a two-stage annealing in which the temperature is kept below the Ar 1 transformation point for 20 hours or more.
　以下、本発明の高炭素熱延鋼板の製造方法における限定理由について説明する。なお、説明において、温度に関する「℃」表示は、鋼板表面あるいは鋼素材の表面における温度を表すものとする。 The reasons for limitation in the method for producing a high carbon hot rolled steel sheet according to the present invention will be described below. In the description, “° C.” regarding temperature indicates the temperature on the surface of the steel plate or the surface of the steel material.
　本発明において、鋼素材の製造方法は、特に限定する必要はない。例えば、本発明の高炭素鋼を溶製するには、転炉、電気炉どちらも使用可能である。転炉等の公知の方法で溶製された高炭素鋼は、造塊－分塊圧延または連続鋳造によりスラブ等（鋼素材）とされる。スラブは、通常、加熱された後、熱間圧延（熱間粗圧延、仕上圧延）される。 In the present invention, the manufacturing method of the steel material does not need to be particularly limited. For example, both a converter and an electric furnace can be used to produce the high carbon steel of the present invention. High carbon steel melted by a known method such as a converter is made into a slab (steel material) by ingot-bulk rolling or continuous casting. The slab is usually heated and then hot-rolled (hot rough rolling, finish rolling).
　例えば、連続鋳造で製造されたスラブの場合は、そのままあるいは温度低下を抑制する目的で保熱して、圧延する直送圧延を適用してもよい。また、スラブを加熱して熱間圧延する場合は、スケールによる表面状態の劣化を避けるために、スラブの加熱温度を１２８０℃以下とすることが好ましい。なお、スラブの加熱温度の下限については１１００℃が好ましく、１１５０℃がさらに好ましく、１２００℃以上がより一層好ましい。なお、熱間圧延では、仕上圧延終了温度を確保するため、熱間圧延中にシートバーヒータ等の加熱手段により被圧延材の加熱を行ってもよい。 For example, in the case of a slab manufactured by continuous casting, direct feed rolling may be applied as it is or with heat retention for the purpose of suppressing the temperature decrease and rolling. Further, when the slab is heated and hot-rolled, the heating temperature of the slab is preferably 1280° C. or lower in order to avoid deterioration of the surface state due to scale. The lower limit of the heating temperature of the slab is preferably 1100°C, more preferably 1150°C, and even more preferably 1200°C or higher. In the hot rolling, in order to secure the finish rolling finish temperature, the material to be rolled may be heated by a heating means such as a sheet bar heater during the hot rolling.
　仕上圧延終了温度：Ａｒ３変態点以上で仕上圧延
　仕上圧延終了温度がＡｒ３変態点未満では、熱間圧延後および焼鈍後に粗大なフェライト粒が形成され、伸びが著しく低下する。このため、仕上圧延終了温度は、Ａｒ３変態点以上とする。好ましくは（Ａｒ３変態点＋２０℃）以上とする。なお、仕上圧延終了温度の上限は、特に規定する必要はないが、仕上圧延後の冷却を円滑に行うためには、１０００℃以下とすることが好ましい。 Finishing rolling end temperature: Finish rolling at Ar 3 transformation point or higher If the finishing rolling termination temperature is less than Ar 3 transformation point, coarse ferrite grains are formed after hot rolling and after annealing, and elongation is remarkably reduced. Therefore, the finish rolling end temperature is set to the Ar 3 transformation point or higher. The temperature is preferably (Ar 3 transformation point+20° C.) or higher. The upper limit of the finish rolling finish temperature is not particularly limited, but it is preferably 1000° C. or lower for smooth cooling after finish rolling.
　なお、上述したＡｒ３変態点は、フォーマスター試験などによる冷却時の熱膨張測定や電気抵抗測定による実測により決定することができる。 The Ar 3 transformation point described above can be determined by the thermal expansion measurement during cooling such as the Formaster test or the actual measurement by the electric resistance measurement.
　仕上圧延後、平均冷却速度：２０～１００℃／ｓｅｃで６５０～７５０℃まで冷却
　仕上圧延後、６５０～７５０℃までの平均冷却速度は焼鈍後の球状化セメンタイトのサイズに大きく影響する。仕上圧延後、平均冷却速度が２０℃／ｓｅｃ未満では、焼鈍前組織としてフェライト組織が多すぎるフェライトとパーライト組織になるため、焼鈍後所定のセメンタイト分散状態やサイズが得られない。そのため、２０℃／ｓｅｃ以上で冷却する必要がある。好ましくは２５℃／ｓｅｃ以上である。一方、平均冷却速度が１００℃／ｓｅｃを超えると焼鈍後に所定のサイズを有するセメンタイトが得られにくくなるため、１００℃／ｓｅｃ以下とする。好ましくは７５℃／ｓｅｃ以下である。 After finish rolling, the average cooling rate is 20 to 100° C./sec and is cooled to 650 to 750° C. After finish rolling, the average cooling rate from 650 to 750° C. greatly affects the size of spheroidized cementite after annealing. If the average cooling rate after finish rolling is less than 20° C./sec, the ferrite structure and the pearlite structure are too large as the pre-annealing structure, so that the predetermined cementite dispersed state and size cannot be obtained after the annealing. Therefore, it is necessary to cool at 20° C./sec or more. It is preferably 25° C./sec or more. On the other hand, if the average cooling rate exceeds 100° C./sec, it becomes difficult to obtain cementite having a predetermined size after annealing, so it is set to 100° C./sec or less. It is preferably 75° C./sec or less.
　巻取温度：５００～７００℃
　仕上圧延後の熱延鋼板は、コイル形状に巻き取られる。巻取温度が高すぎると熱延鋼板の強度が低くなり過ぎて、コイル形状に巻き取られた際、コイルの自重で変形する場合がある。このため、操業上の観点から好ましくない。したがって、巻取温度の上限を７００℃とする。好ましくは６９０℃以下である。一方、巻取温度が低すぎると熱延鋼板が硬質化するため、好ましくない。したがって、巻取温度は５００℃とする。好ましくは５３０℃以上である。 Winding temperature: 500-700°C
The hot rolled steel sheet after finish rolling is wound into a coil shape. If the coiling temperature is too high, the strength of the hot-rolled steel sheet becomes too low, and when coiled into a coil shape, the coil may be deformed by its own weight. Therefore, it is not preferable from the viewpoint of operation. Therefore, the upper limit of the winding temperature is 700°C. The temperature is preferably 690°C or lower. On the other hand, if the winding temperature is too low, the hot-rolled steel sheet becomes hard, which is not preferable. Therefore, the winding temperature is 500°C. It is preferably 530° C. or higher.
　コイル状に巻き取った後、常温まで冷却し、酸洗処理を施しても良い。酸洗処理後、焼鈍を行う。なお、酸洗処理は公知の方法を適用できる。その後、得られた熱延鋼板に以下の焼鈍を施す。 After coiling, it may be cooled to room temperature and pickled. After the pickling treatment, annealing is performed. A known method can be applied to the pickling treatment. Then, the obtained hot rolled steel sheet is annealed as follows.
　４５０～６００℃の温度範囲の平均加熱速度：１５℃／ｈ以上
　上記のようにして得た熱延鋼板に、焼鈍（セメンタイトの球状化焼鈍）を施す。窒素雰囲気中での焼鈍では、４５０～６００℃の温度範囲ではアンモニアガスが発生しやすくなり、アンモニアガスから分解された窒素が表面鋼板に入り、鋼中のＢやＡｌと結合し窒化物を生成する。そのため、４５０～６００℃の温度範囲の加熱時間はできるだけ短くする。この温度範囲での平均加熱速度は、１５℃／ｈ以上とする。生産性向上を目的として炉内ばらつきを抑制する観点から、好ましくは１００℃／ｈ以下とし、さらに好ましくは７０℃／ｈ以下とする。 Average heating rate in the temperature range of 450 to 600° C.: 15° C./h or more The hot rolled steel sheet obtained as described above is annealed (cementite spheroidizing annealing). When annealed in a nitrogen atmosphere, ammonia gas is likely to be generated in the temperature range of 450 to 600° C., and nitrogen decomposed from the ammonia gas enters the surface steel sheet and combines with B and Al in the steel to form a nitride. To do. Therefore, the heating time in the temperature range of 450 to 600° C. should be as short as possible. The average heating rate in this temperature range is 15° C./h or more. From the viewpoint of suppressing variations in the furnace for the purpose of improving productivity, it is preferably 100° C./h or less, more preferably 70° C./h or less.
　焼鈍温度：Ａｃ１変態点未満で１．０ｈ以上保持
　焼鈍温度がＡｃ１変態点以上であると、オーステナイトが析出し、焼鈍後の冷却過程において粗大なパーライト組織が形成され、不均一な組織となる。このため、焼鈍温度は、Ａｃ１変態点未満とする。好ましくは（Ａｃ１変態点－１０℃）以下である。なお、焼鈍温度の下限は特に定めないが、所定のセメンタイト分散状態を得るには、焼鈍温度は６００℃以上が好ましく、さらに好ましくは７００℃以上である。なお、雰囲気ガスは、窒素、水素、窒素と水素の混合ガスのいずれも使用できる。また、上記焼鈍温度における保持時間は、１．０時間（ｈ）以上とする。焼鈍温度における保持時間が１．０時間未満であると、焼鈍の効果が乏しく、本発明の目標とする組織が得られず、その結果、本発明の目標とする鋼板の硬さおよび伸びが得られない。したがって、焼鈍温度における保持時間は１．０時間以上とする。好ましくは５時間以上であり、さらに好ましくは２０時間超えである。一方、上記焼鈍温度における保持時間が４０．０時間を超えると、生産性が低下し、製造コストが過大となる。そのため、上記焼鈍温度における保持時間は、４０．０時間以下とすることが好ましい。さらに好ましくは３５時間以下である。 Annealing temperature: Hold for 1.0 h or more below Ac 1 transformation point If the annealing temperature is Ac 1 transformation point or more, austenite precipitates and a coarse pearlite structure is formed in the cooling process after annealing, resulting in a non-uniform structure. Become. Therefore, the annealing temperature is lower than the Ac 1 transformation point. It is preferably (Ac 1 transformation point −10° C.) or less. Although the lower limit of the annealing temperature is not particularly defined, the annealing temperature is preferably 600° C. or higher, more preferably 700° C. or higher in order to obtain a predetermined cementite dispersed state. The atmosphere gas may be nitrogen, hydrogen, or a mixed gas of nitrogen and hydrogen. The holding time at the annealing temperature is 1.0 hour (h) or more. If the holding time at the annealing temperature is less than 1.0 hour, the effect of annealing is poor, and the target structure of the present invention cannot be obtained. As a result, the hardness and elongation of the steel plate targeted by the present invention are obtained. I can't. Therefore, the holding time at the annealing temperature is set to 1.0 hour or more. It is preferably 5 hours or more, more preferably 20 hours or more. On the other hand, if the holding time at the annealing temperature exceeds 40.0 hours, the productivity will decrease and the manufacturing cost will be excessive. Therefore, the holding time at the annealing temperature is preferably 40.0 hours or less. More preferably, it is 35 hours or less.
　本発明では、上記した焼鈍に代えて以下の２段焼鈍を施すことができる。具体的には、巻き取り、常温まで冷却して熱延鋼板とした後、この熱延鋼板を平均加熱速度：１５℃／ｈ以上で４５０～６００℃の温度範囲を加熱し、Ａｃ１変態点以上Ａｃ３変態点以下で０．５ｈ以上保持（１段目の焼鈍）し、次いで平均冷却速度：１～２０℃／ｈでＡｒ１変態点未満に冷却し、Ａｒ１変態点未満で２０ｈ以上保持（２段目の焼鈍）する２段焼鈍を施すことにより製造することも可能である。 In the present invention, the following two-step annealing can be applied instead of the above-mentioned annealing. Specifically, after winding and cooling to room temperature to form a hot-rolled steel sheet, the hot-rolled steel sheet is heated in a temperature range of 450 to 600° C. at an average heating rate of 15° C./h or more to obtain an Ac 1 transformation point. Hold at 0.5 h or more below the Ac 3 transformation point (first annealing), then cool to less than Ar 1 transformation point at average cooling rate: 1 to 20° C./h, and 20 h or more below Ar 1 transformation point It is also possible to manufacture by performing a two-step annealing that holds (second-step annealing).
　本発明では、熱延鋼板を平均加熱速度：１５℃／ｈ以上で４５０～６００℃の温度範囲を加熱し、Ａｃ１変態点以上で０．５ｈ以上保持し、熱延鋼板中に析出していた比較的微細な炭化物を溶解してγ相中に固溶させ、その後、平均冷却速度：１～２０℃／ｈでＡｒ１変態点未満に冷却し、Ａｒ１変態点未満で２０ｈ以上保持する。これより、比較的粗大な未溶解炭化物等を核として固溶Ｃを析出させて、全体のセメンタイト数に対する円相当直径０．１μｍ以下のセメンタイト数の割合が２０％以下となるような、炭化物（セメンタイト）の分散を制御された状態とすることができる。すなわち、本発明では、所定条件で２段焼鈍を施すことで、炭化物の分散形態を制御し、鋼板を軟質化させる。本発明で対象とする高炭素鋼板では、軟質化する上で焼鈍後における炭化物の分散形態を制御することが重要となる。本発明では、高炭素熱延鋼板をＡｃ１変態点以上Ａｃ３変態点以下で保持する（１段目の焼鈍）ことで、微細な炭化物を溶解するとともに、Ｃをγ（オーステナイト）中に固溶する。その後のＡｒ１変態点未満の冷却段階や保持段階（２段目の焼鈍）において、Ａｃ１変態点以上の温度域で存在するα／γ界面や未溶解炭化物が核生成サイトとなり、比較的粗大な炭化物が析出する。以下、このような２段焼鈍の条件について説明する。なお、焼鈍の際の雰囲気ガスは、窒素、水素、窒素と水素の混合ガスのいずれも使用できる。 In the present invention, the hot-rolled steel sheet is heated in the temperature range of 450 to 600° C. at an average heating rate of 15° C./h or more and kept at the Ac 1 transformation point or more for 0.5 h or more to precipitate in the hot-rolled steel sheet. The relatively fine carbide is dissolved to form a solid solution in the γ phase, and thereafter, the average cooling rate is cooled to below Ar 1 transformation point at an average cooling rate of 1 to 20° C./h, and maintained for 20 hours or more below Ar 1 transformation point. .. From this, a solid solution C is precipitated by using a relatively coarse undissolved carbide as a nucleus, and the ratio of the number of cementites having a circle-equivalent diameter of 0.1 μm or less to the total number of cementites is 20% or less ( Cementite) dispersion can be controlled. That is, in the present invention, the two-step annealing is performed under predetermined conditions to control the dispersed form of the carbide and soften the steel sheet. In the high carbon steel sheet targeted by the present invention, it is important to control the dispersed form of carbides after annealing in order to soften the steel. In the present invention, by holding the high carbon hot-rolled steel sheet at the Ac 1 transformation point or more and the Ac 3 transformation point or less (first-stage annealing), fine carbides are dissolved and C is solidified in γ (austenite). Melt. In the subsequent cooling step below the Ar 1 transformation point and the holding step (second annealing), the α/γ interface and undissolved carbides existing in the temperature range above the Ac 1 transformation point become nucleation sites and are relatively coarse. Carbide precipitates. Hereinafter, the conditions of such two-step annealing will be described. As the atmosphere gas at the time of annealing, any of nitrogen, hydrogen, and a mixed gas of nitrogen and hydrogen can be used.
　４５０～６００℃の温度範囲の平均加熱速度：１５℃／ｈ以上
　上記と同じ理由で、４５０～６００℃の温度範囲ではアンモニアガスが発生しやすくなり、アンモニアガスから分解された窒素が表面鋼板に入り、鋼中のＢやＡｌと結合し窒化物を生成するため、４５０～６００℃の温度範囲の加熱時間はできるだけ短くする。この温度範囲での平均加熱速度は、１５℃／ｈ以上とする。好ましくは２０℃／ｈ以上とする。平均加熱速度の上限は１００℃／ｈとすることが好ましく、さらに好ましくは９０℃／ｈ以下とする。 Average heating rate in the temperature range of 450 to 600° C.: 15° C./h or more For the same reason as above, ammonia gas is easily generated in the temperature range of 450 to 600° C., and nitrogen decomposed from the ammonia gas becomes surface steel sheet. The heating time in the temperature range of 450 to 600° C. is made as short as possible because it enters and combines with B and Al in the steel to form a nitride. The average heating rate in this temperature range is 15° C./h or more. It is preferably 20° C./h or more. The upper limit of the average heating rate is preferably 100°C/h, more preferably 90°C/h or less.
　Ａｃ１変態点以上Ａｃ３変態点以下で０．５ｈ以上保持（１段目の焼鈍）
　熱延鋼板をＡｃ１変態点以上で保持することにより、鋼板組織のフェライトの一部をオーステナイトに変態させ、フェライト中に析出していた微細な炭化物を溶解させ、Ｃをオーステナイト中に固溶させる。一方、オーステナイトに変態せずに残ったフェライトは高温で焼鈍されるため、転位密度が減少して軟化する。また、フェライト中には溶解しなかった比較的粗大な炭化物（未溶解炭化物）が残存するが、オストワルド成長により、より粗大になる。焼鈍温度がＡｃ１変態点未満では、オーステナイト変態が生じないため、炭化物をオーステナイト中に固溶させることができない。一方、１段目の焼鈍温度がＡｃ３変態点超になると焼鈍後に棒状のセメンタイトが多数得られて所定の伸びが得られないため、Ａｃ３変態点以下とする。また、本発明では、Ａｃ１変態点以上Ａｃ３変態点以下での保持時間が０．５ｈ未満では微細な炭化物を十分に溶解することができない。このため、１段目の焼鈍として、Ａｃ１変態点以上Ａｃ３変態点以下で０．５ｈ以上保持することとする。保持時間は、好ましくは１．０ｈ以上とする。また、保持時間は１０ｈ以下とすることが好ましい。なお、Ａｃ１変態点以上Ａｃ３変態点以下で保持して焼鈍を行う場合でも、加熱速度は、４５０～６００℃の温度範囲の平均加熱速度を１５℃／ｈ以上とし、上限を１００℃/ｈ以下とすることが好ましい。 Hold for 0.5 h or more above the Ac 1 transformation point and below the Ac 3 transformation point (first-stage annealing)
By holding the hot-rolled steel sheet at the Ac 1 transformation point or higher, a part of the ferrite of the steel sheet structure is transformed into austenite, the fine carbides precipitated in the ferrite are dissolved, and C is dissolved in austenite. .. On the other hand, the ferrite remaining without being transformed into austenite is annealed at a high temperature, so that the dislocation density decreases and the ferrite softens. Further, although relatively coarse carbides (undissolved carbides) that have not been dissolved remain in the ferrite, they become coarser due to Ostwald growth. If the annealing temperature is lower than the Ac 1 transformation point, austenite transformation does not occur, so that the carbide cannot be dissolved in austenite. On the other hand, when the annealing temperature of the first step exceeds the Ac 3 transformation point, a large number of rod-shaped cementites are obtained after annealing and a predetermined elongation cannot be obtained, so the temperature is set to the Ac 3 transformation point or lower. Further, in the present invention, if the holding time at the Ac 1 transformation point or more and the Ac 3 transformation point or less is less than 0.5 h, fine carbides cannot be sufficiently dissolved. For this reason, as the first-stage annealing, 0.5 h or more is maintained at the Ac 1 transformation point or more and the Ac 3 transformation point or less. The holding time is preferably 1.0 h or longer. The holding time is preferably 10 hours or less. Even when annealing is performed while maintaining the temperature at the Ac 1 transformation point or more and the Ac 3 transformation point or less, the heating rate is 15°C/h or more in the average heating rate in the temperature range of 450 to 600°C, and the upper limit is 100°C/ It is preferably h or less.
　平均冷却速度：１～２０℃／ｈでＡｒ１変態点未満に冷却
　上記した１段目の焼鈍の後、２段目の焼鈍の温度域であるＡｒ１変態点未満に、平均冷却速度：１～２０℃／ｈで冷却する。冷却途中に、オーステナイトからフェライトへの変態に伴いオーステナイトから吐き出されるＣが、α／γ界面や未溶解炭化物を核生成サイトとして、比較的粗大な球状炭化物として析出する。この冷却においては、パーライトが生成しないように冷却速度を調整する必要がある。１段目の焼鈍後、２段目の焼鈍までの平均冷却速度が、１℃／ｈ未満では生産効率が悪いため、該平均冷却速度は１℃／ｈ以上とする。好ましくは５℃／ｈ以上とする。一方、平均冷却速度が２０℃／ｈを超えて大きくなると、パーライトが析出し、硬度が高くなるため、２０℃／ｈ以下とする。好ましくは１５℃／ｈ以下とする。 Average cooling rate: cooling to below Ar 1 transformation point at 1 to 20° C./h After the above first annealing step, below the Ar 1 transformation point which is the temperature range of the second annealing step, average cooling rate: 1 Cool at ~20°C/h. During the cooling, C discharged from the austenite along with the transformation from austenite to ferrite is precipitated as a relatively coarse spherical carbide by using the α/γ interface and undissolved carbide as a nucleation site. In this cooling, it is necessary to adjust the cooling rate so that pearlite is not generated. If the average cooling rate from the first annealing to the second annealing is less than 1° C./h, the production efficiency is poor, so the average cooling rate is 1° C./h or more. It is preferably 5° C./h or more. On the other hand, when the average cooling rate exceeds 20° C./h and becomes large, pearlite precipitates and the hardness increases, so the rate is set to 20° C./h or less. The rate is preferably 15° C./h or less.
　Ａｒ１変態点未満で２０ｈ以上保持（２段目の焼鈍）
　上記した１段目の焼鈍後、所定の平均冷却速度で冷却してＡｒ１変態点未満で保持することで、オストワルド成長により、粗大な球状炭化物をさらに成長させ、微細な炭化物を消失させる。Ａｒ１変態点未満での保持時間が２０ｈ未満では、炭化物を十分に成長させることができず、焼鈍後の硬度が大きくなりすぎる。このため、２段目の焼鈍はＡｒ１変態点未満で２０ｈ以上保持とする。なお、特に限定するものではないが、２段目の焼鈍温度は炭化物を十分成長させるため、６６０℃以上とすることが好ましく、また、保持時間は生産効率の観点から、３０ｈ以下とすることが好ましい。 Hold for 20 h or more below Ar 1 transformation point (second-stage annealing)
After the annealing in the first step described above, cooling is performed at a predetermined average cooling rate and the temperature is maintained below the Ar 1 transformation point, whereby coarse spherical carbides are further grown and fine carbides disappear by Ostwald ripening. If the holding time below the Ar 1 transformation point is less than 20 h, the carbide cannot be grown sufficiently and the hardness after annealing becomes too large. Therefore, the second annealing is held for 20 hours or more below the Ar 1 transformation point. Although not particularly limited, the second annealing temperature is preferably 660° C. or higher in order to sufficiently grow the carbide, and the holding time is 30 h or less from the viewpoint of production efficiency. preferable.
　なお、上述したＡｃ３変態点、Ａｃ１変態点、Ａｒ３変態点、Ａｒ１変態点は、フォーマスター試験などによる加熱時、冷却時の熱膨張測定や電気抵抗測定による実測により決定することができる。 The above Ac 3 transformation point, Ac 1 transformation point, Ar 3 transformation point, and Ar 1 transformation point may be determined by actual measurement by thermal expansion measurement or electric resistance measurement during heating or cooling by the Formaster test or the like. it can.
　また、上述した平均加熱速度、平均冷却速度は、炉内に設置した熱電対で温度を測定し求める。 Also, the above average heating rate and average cooling rate are obtained by measuring the temperature with a thermocouple installed in the furnace.
　表１に示す鋼番Ａ～Ｔの成分組成を有する鋼を溶製し、次いで表２－１および表３－１に示す製造条件に従って、熱間圧延を行った。次いで、酸洗し、窒素雰囲気中（雰囲気ガス：窒素）で、表２－１および表３－１に示す焼鈍温度および焼鈍時間（ｈ）にて焼鈍（球状化焼鈍）を施して、板厚３．０ｍｍの熱延焼鈍板を製造した。 Steels having the component compositions of steel Nos. A to T shown in Table 1 were melted, and then hot rolled according to the manufacturing conditions shown in Table 2-1 and Table 3-1. Then, pickling, and annealing (spheroidizing annealing) in a nitrogen atmosphere (atmosphere gas: nitrogen) at an annealing temperature and an annealing time (h) shown in Tables 2-1 and 3-1 to obtain a sheet thickness A 3.0 mm hot rolled annealed plate was manufactured.
　本発明の実施例では、このようにして得られた熱延焼鈍板から試験片を採取し、下記のように、ミクロ組織、固溶Ｂ量、ＡｌＮ中のＮ量、引張強度、全伸びおよび焼入れ硬さ（焼入れ後の鋼板硬さ、浸炭焼入れ後の鋼板硬さ）をそれぞれ求めた。なお、表１に示すＡｃ３変態点、Ａｃ１変態点、Ａｒ１変態点およびＡｒ３変態点はフォーマスター試験により求めたものである。 In the examples of the present invention, test pieces were sampled from the hot rolled annealed sheet thus obtained, and the microstructure, the amount of solid solution B, the amount of N in AlN, the tensile strength, the total elongation and Hardening hardness (steel plate hardness after quenching, steel plate hardness after carburizing and quenching) was determined. The Ac 3 transformation point, the Ac 1 transformation point, the Ar 1 transformation point and the Ar 3 transformation point shown in Table 1 were obtained by the Formaster test.
　（１）ミクロ組織
　焼鈍後の鋼板のミクロ組織は、板幅中央部から採取した試験片（大きさ：３ｍｍｔ×１０ｍｍ×１０ｍｍ）を切断研磨後、ナイタール腐食を施し、走査型電子顕微鏡（ＳＥＭ）を用いて、表層から板厚１／４のところの５箇所で３０００倍の倍率で撮影した。撮影した組織写真を画像処理により各相（フェライト、セメンタイト、パーライトなど）を特定した。表２－２および表３－２にはミクロ組織として「パーライト面積率」を記載しており、パーライトが面積率で６．５％を超えて認められた鋼については、比較例としている。面積率で６．５％以下のパーライトと、フェライトと、セメンタイトを有する鋼については、本発明例としている。 (1) Microstructure The microstructure of the annealed steel plate was obtained by cutting and polishing a test piece (size: 3 mmt x 10 mm x 10 mm) taken from the center of the plate width, and then subjecting it to nital corrosion, and scanning electron microscope (SEM). The images were taken at a magnification of 3000 times at 5 positions from the surface layer at a plate thickness of 1/4. Each phase (ferrite, cementite, pearlite, etc.) was specified by image processing of the photographed microstructure. In Tables 2-2 and 3-2, the “perlite area ratio” is described as a microstructure, and the steel in which pearlite is found to exceed 6.5% in area ratio is taken as a comparative example. Steels having a pearlite area ratio of 6.5% or less, ferrite, and cementite are examples of the present invention.
　また、ＳＥＭ画像から画像解析ソフトを用いて、フェライトとフェライト以外の領域とを二値化して、フェライトの面積率（％）を求めた。セメンタイトも同様に、ＳＥＭ画像から画像解析ソフトを用いて、セメンタイトとセメンタイト以外の領域とを二値化して、セメンタイトの面積率（％）を求めた。また、パーライトは、１００（％）からフェライトとセメンタイトの各面積率（％）を引いた値を、パーライトの面積率（％）とした。 Also, the area ratio (%) of the ferrite was obtained by binarizing the ferrite and the area other than the ferrite using image analysis software from the SEM image. Similarly, for cementite, the area ratio (%) of cementite was obtained by binarizing the cementite and the region other than the cementite from the SEM image using image analysis software. For pearlite, the value obtained by subtracting the area ratio (%) of each of ferrite and cementite from 100 (%) was defined as the area ratio (%) of pearlite.
　また、撮影した組織写真について、個々のセメンタイト径を評価した。セメンタイト径は、長径と短径を測定し、円相当直径に換算した。平均セメンタイト径は円相当直径に換算した全てのセメンタイトの円相当直径の合計をセメンタイト総数で除して求めた。円相当直径の値が０．１μｍ以下のセメンタイトの個数を測定し、円相当直径０．１μｍ以下のセメンタイトの数とした。また、全セメンタイトの個数を求め、全セメンタイト数とした。そして、全セメンタイト数に対する円相当直径０．１μｍ以下のセメンタイト数の割合（（円相当直径０．１μｍ以下のセメンタイト数／全セメンタイト数）×１００（％））を求めた。なお、この「円相当直径０．１μｍ以下のセメンタイトの割合」を、円相当直径０．１μｍ以下のセメンタイトと単に称する場合もある。 Also, individual cementite diameters were evaluated for the taken micrographs. For the cementite diameter, the major axis and the minor axis were measured and converted into a circle equivalent diameter. The average cementite diameter was calculated by dividing the sum of the equivalent circle diameters of all the cementites converted to the equivalent circle diameter by the total number of cementites. The number of cementites having a circle equivalent diameter of 0.1 μm or less was measured and used as the number of cementite having a circle equivalent diameter of 0.1 μm or less. In addition, the total number of cementites was calculated and used as the total number of cementites. Then, the ratio of the number of cementites having a circle-equivalent diameter of 0.1 μm or less to the total number of cementites ((the number of cementites having a circle-equivalent diameter of 0.1 μm or less/total number of cementites)×100(%)) was determined. The "ratio of cementite having a circle-equivalent diameter of 0.1 µm or less" may be simply referred to as cementite having a circle-equivalent diameter of 0.1 µm or less.
　また、撮影した組織写真について、ＪＩＳ　Ｇ　０５５１に定められた結晶粒度の評価方法（切断法）を用いて、フェライトの平均粒径を求めた。 Also, the average grain size of ferrite was determined for the photographed structure using the grain size evaluation method (cutting method) specified in JIS G 0551.
　（２）固溶Ｂ量の平均濃度の測定
　下記参考文献に記載されている方法と同じ手法で求めた。すなわち、表層から深さ１００μｍまでの領域の研削粉を収集して３回測定し、この平均値を固溶Ｂ量の平均濃度として求めた。
［参考文献］城代哲史、石田智治、猪瀬国生、藤本京子,鉄と鋼,vol.99 (2013) No.5, p.362-365
　（３）ＡｌＮとして存在するＮ量の平均濃度の測定
　上記と同様、下記参考文献に記載されている方法と同じ手法で、ＡｌＮとして存在するＮ量の平均濃度を求めた。
［参考文献］城代哲史、石田智治、猪瀬国生、藤本京子,鉄と鋼,vol.99(2013) No.5, p.362-365
　（４）鋼板の引張強度と伸び
　焼鈍後の鋼板（原板）から、圧延方向に対して０°の方向（Ｌ方向）に切り出したＪＩＳ５号引張試験片を用いて、１０ｍｍ／分で引張試験を行い、公称応力公称歪曲線を求め、最大応力を引張強度とした。また、破断したサンプルを突き合わせて全伸びを求めた。その結果を、伸び（Ｅｌ）とした。 (2) Measurement of average concentration of solid solution B amount It was determined by the same method as described in the following references. That is, grinding powder in a region from the surface layer to a depth of 100 μm was collected and measured three times, and the average value was determined as the average concentration of the solid solution B amount.
[Reference] Satoshi Joshiro, Tomoji Ishida, Kunio Inose, Kyoko Fujimoto, Iron and Steel, vol.99 (2013) No.5, p.362-365
(3) Measurement of average concentration of N amount present as AlN Similarly to the above, the average concentration of N amount present as AlN was determined by the same method as described in the following references.
[Reference] Satoshi Joshiro, Tomoji Ishida, Kunio Inose, Kyoko Fujimoto, Iron and Steel, vol.99 (2013) No.5, p.362-365
(4) Tensile Strength and Elongation of Steel Plate A tensile test was performed at 10 mm/min using a JIS No. 5 tensile test piece cut out from the annealed steel plate (original plate) in the direction of 0° (L direction) with respect to the rolling direction. Then, the nominal stress and nominal strain curve were obtained, and the maximum stress was taken as the tensile strength. Further, the broken samples were butted against each other to determine the total elongation. The result was defined as elongation (El).
　（５）焼入れ後の鋼板硬さ（ズブ焼入れ性）
　焼鈍後の鋼板の板幅中央から平板試験片（幅１５ｍｍ×長さ４０ｍｍ×板厚３ｍｍ）を採取し、以下のように７０℃油冷により焼入れ処理を施して、焼入れ硬さ（ズブ焼入れ性）を求めた。焼入れ処理は、上記平板試験片を用いて９００℃で６００ｓ保持して直ちに７０℃の油で冷却する方法（７０℃油冷）で実施した。焼入れ硬さは、焼入れ処理後の試験片の切断面について、表層から７０μｍ板厚内部の領域と１／４板厚にてビッカース硬さ試験機で荷重０．２ｋｇｆの条件下で、硬さを５点測定し、平均硬さを求め、これを焼入れ硬さ（ＨＶ）とした。なお、上記した表層から７０μｍ板厚内部の領域は、表２－２および表３－２において「表層」と示す。 (5) Steel plate hardness after quenching (Zub hardenability)
A flat plate test piece (width 15 mm × length 40 mm × plate thickness 3 mm) was taken from the center of the plate width of the annealed steel plate and subjected to quenching treatment by oil cooling at 70°C as described below to obtain quenching hardness (dub quenchability). ) Was asked. The quenching treatment was carried out by using the above flat plate test piece and holding it at 900° C. for 600 s and immediately cooling with oil at 70° C. (70° C. oil cooling). The quenching hardness is the hardness of the cut surface of the test piece after quenching under the condition of a load of 0.2 kgf with a Vickers hardness tester in an area within the thickness of 70 μm from the surface layer and a quarter thickness. Five points were measured and the average hardness was determined, which was taken as the quenching hardness (HV). The region within the plate thickness of 70 μm from the surface layer is indicated as “surface layer” in Table 2-2 and Table 3-2.
　（６）浸炭焼入れ後の鋼板硬さ（浸炭焼入れ性）
　焼鈍後の鋼板について、９３０℃で鋼の均熱、浸炭処理、拡散処理といった浸炭焼入れ処理を合計時間４時間で行い、８５０℃で３０分保持した後、油冷した（油冷の温度：６０℃）。鋼板表面からの深さ０．１ｍｍの位置と深さ１．２ｍｍの位置まで０．１ｍｍ間隔にて硬さを荷重１ｋｇｆの条件下で測定し、浸炭焼入れ時の表層０．１ｍｍの硬さ（ＨＶ）と有効硬化層深さ（ｍｍ）を求めた。有効硬化層深さとは、熱処理後表面から硬さを測定し、５５０ＨＶ以上となる深さと定義する。 (6) Steel plate hardness after carburizing and quenching (carburizing and quenching property)
The annealed steel sheet was subjected to carburizing and quenching treatment such as soaking, carburizing treatment, and diffusion treatment at 930° C. for a total time of 4 hours, and was held at 850° C. for 30 minutes and then oil cooled (oil cooling temperature: 60 C). The hardness was measured at a depth of 0.1 mm from the surface of the steel sheet at a depth of 1.2 mm at 0.1 mm intervals under a load of 1 kgf, and the hardness of the surface layer at the time of carburizing and quenching was 0.1 mm ( HV) and effective hardened layer depth (mm) were determined. The effective hardened layer depth is defined as the depth at which 550 HV or more is obtained by measuring the hardness from the surface after heat treatment.
　そして、上記（５）、（６）より得られた結果から、表４に示す条件で焼入れ性評価を行った。表４は、焼入れ性が十分であると評価できる、Ｃ含有量に応じた焼入れ性の合格規準を表したものである。７０℃油冷後硬さ（ＨＶ）、浸炭焼入れ時の表層０．１ｍｍの深さにおける硬さ（ＨＶ）および浸炭焼入れ時の有効硬化層深さの全てが、表４の規準を満足した場合、合格（記号：○で示す）と判定し、焼入れ性に優れると評価した。一方、いずれかの値が表４に示す規準を満足しない場合、不合格（記号：×で示す）と判定し、焼入れ性に劣ると評価した。 Then, from the results obtained from the above (5) and (6), the hardenability was evaluated under the conditions shown in Table 4. Table 4 shows the acceptance criteria of the hardenability according to the C content, which can be evaluated as having sufficient hardenability. When hardness (HV) after oil cooling at 70°C, hardness (HV) at a depth of 0.1 mm of surface layer during carburizing and quenching, and effective hardened layer depth during carburizing and quenching all satisfy the criteria in Table 4 Was judged to be acceptable (symbol: indicated by ◯) and evaluated to be excellent in hardenability. On the other hand, when any of the values did not satisfy the criteria shown in Table 4, it was judged as unacceptable (symbol: x) and evaluated as being inferior in hardenability.
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000003
Figure JPOXMLDOC01-appb-T000003
Figure JPOXMLDOC01-appb-T000004
Figure JPOXMLDOC01-appb-T000004
Figure JPOXMLDOC01-appb-T000005
Figure JPOXMLDOC01-appb-T000005
Figure JPOXMLDOC01-appb-T000006
Figure JPOXMLDOC01-appb-T000006
　表２－２および表３－２の結果より、本発明例の高炭素熱延鋼板は、全セメンタイト数に対する円相当直径０．１μｍ以下のセメンタイト数の割合が２０％以下であり、平均セメンタイト径が２．５μｍ以下、全ミクロ組織に対する前記セメンタイトの占める割合が３．５％以上１０．０％以下であり、フェライトとセメンタイトを有するミクロ組織を有し、冷間加工性に優れるとともに、焼入れ性にも優れていることがわかる。また、引張強度が４８０ＭＰａ以下、全伸び（Ｅｌ）が３３％以上と優れた機械特性も得ることができた。 From the results of Table 2-2 and Table 3-2, in the high carbon hot-rolled steel sheet of the present invention example, the ratio of the number of cementites having a circle equivalent diameter of 0.1 μm or less to the total number of cementites was 20% or less, and the average cementite diameter was Is 2.5 μm or less, the ratio of the cementite to the total microstructure is 3.5% or more and 10.0% or less, and has a microstructure having ferrite and cementite, and is excellent in cold workability and hardenability. It turns out that it is also excellent. Also, excellent mechanical properties such as tensile strength of 480 MPa or less and total elongation (El) of 33% or more could be obtained.
　一方、本発明の範囲を外れる比較例は、成分組成、ミクロ組織、固溶Ｂ量、ＡｌＮ中のＮ量のいずれか１つ以上が本発明の範囲を満足せず、その結果、冷間加工性、焼入れ性のいずれか１つ以上が、上述の目標性能を満足できないことがわかる。また、引張強度（ＴＳ）、全伸び（Ｅｌ）の１つ以上が目標特性を満足することができないものもあった。例えば、表２－２および表３－２において、鋼ＳはＣ量が本発明範囲よりも低いため、ズブ焼入れ性を満足しない。また、鋼ＴはＣ量が本発明範囲よりも高いため、鋼板のＴＳ、全伸びの特性を満足しない。 On the other hand, in Comparative Examples outside the scope of the present invention, any one or more of the component composition, the microstructure, the amount of solid solution B, and the amount of N in AlN do not satisfy the scope of the present invention, and as a result, cold working It can be seen that any one or more of the hardenability and the hardenability cannot satisfy the above target performance. Further, in some cases, one or more of tensile strength (TS) and total elongation (El) could not satisfy the target characteristics. For example, in Tables 2-2 and 3-2, steel S does not satisfy the zub hardenability because the C content is lower than the range of the present invention. Further, since the steel T has a C content higher than the range of the present invention, it does not satisfy the TS and total elongation characteristics of the steel sheet.

Claims (6)

  1. 　質量％で、
    Ｃ：０．２０％以上０．５０％以下、
    Ｓｉ：０．８％以下、
    Ｍｎ：０．１０％以上０．８０％以下、
    Ｐ：０．０３％以下、
    Ｓ：０．０１０％以下、
    ｓｏｌ．Ａｌ：０．１０％以下、
    Ｎ：０．０１％以下、
    Ｃｒ：１．０％以下、
    Ｂ：０．０００５％以上０．００５％以下、
    さらにＳｂおよびＳｎから選んだ１種または２種を合計で０．００２％以上０．１％以下を含有し、
    残部がＦｅおよび不可避的不純物からなる成分組成を有し、
    ミクロ組織は、
    フェライト、セメンタイト、および全ミクロ組織に対して面積率で６．５％以下の割合を占めるパーライトを有し、
    前記セメンタイトは、全セメンタイト数に対する円相当直径０．１μｍ以下のセメンタイト数の割合が２０％以下、平均セメンタイト径が２．５μｍ以下、全ミクロ組織に対する前記セメンタイトの占める割合が面積率で３．５％以上１０．０％以下であり、
    表層から深さ１００μｍまでの領域における固溶Ｂ量の平均濃度が１０質量ｐｐｍ以上であり、
    表層から深さ１００μｍまでの領域におけるＡｌＮとして存在するＮ量の平均濃度が７０質量ｐｐｍ以下である高炭素熱延鋼板。     In mass %,
    C: 0.20% or more and 0.50% or less,
    Si: 0.8% or less,
    Mn: 0.10% or more and 0.80% or less,
    P: 0.03% or less,
    S: 0.010% or less,
    sol. Al: 0.10% or less,
    N: 0.01% or less,
    Cr: 1.0% or less,
    B: 0.0005% or more and 0.005% or less,
    Further, one or two kinds selected from Sb and Sn are contained in a total amount of 0.002% or more and 0.1% or less,
    The balance has a composition of Fe and inevitable impurities,
    The microstructure is
    Ferrite, cementite, and pearlite occupying 6.5% or less in area ratio with respect to the entire microstructure,
    In the cementite, the ratio of the number of cementites having a circle-equivalent diameter of 0.1 μm or less to the total number of cementites is 20% or less, the average cementite diameter is 2.5 μm or less, and the ratio of the cementite to the entire microstructure is 3.5 in area ratio. % Or more and 10.0% or less,
    The average concentration of solid solution B in the region from the surface layer to a depth of 100 μm is 10 mass ppm or more,
    A high carbon hot-rolled steel sheet having an average concentration of 70 mass ppm or less of N existing as AlN in a region from the surface layer to a depth of 100 μm.
  2. 　引張強度が４８０ＭＰａ以下、全伸びが３３％以上である請求項１に記載の高炭素熱延鋼板。 The high carbon hot rolled steel sheet according to claim 1, which has a tensile strength of 480 MPa or less and a total elongation of 33% or more.
  3. 　前記フェライトの平均粒径が４～２５μｍである請求項１または２に記載の高炭素熱延鋼板。 The high carbon hot-rolled steel sheet according to claim 1 or 2, wherein the ferrite has an average particle size of 4 to 25 µm.
  4. 　前記成分組成に加えてさらに、質量％で、下記Ａ群およびＢ群のうちから選ばれた１群または２群を含有する請求項１～３のいずれかに記載の高炭素熱延鋼板。
    　　　　　　　　　　　　　　　　記
    Ａ群：Ｔｉ：０．０６％以下
    Ｂ群：Ｎｂ、Ｍｏ、Ｔａ、Ｎｉ、Ｃｕ、Ｖ、Ｗのうちから選ばれた１種または２種以上を、それぞれ０．０００５％以上０．１％以下     The high carbon hot-rolled steel sheet according to any one of claims 1 to 3, which further comprises, in mass%, one or two groups selected from the following group A and group B in addition to the component composition.
    Note Group A: Ti: 0.06% or less Group B: One or two or more selected from Nb, Mo, Ta, Ni, Cu, V, and W, respectively, 0.0005% or more 0.1 %Less than
  5. 　請求項１～４のいずれかに記載の高炭素熱延鋼板の製造方法であって、
    前記成分組成を有する鋼を、熱間粗圧延後、仕上圧延終了温度：Ａｒ３変態点以上で仕上圧延を行い、その後、平均冷却速度：２０～１００℃／ｓｅｃで６５０～７５０℃まで冷却し、
    巻取温度：５００～７００℃で巻き取り、熱延鋼板とした後、
    該熱延鋼板を、平均加熱速度：１５℃／ｈ以上で４５０～６００℃の温度範囲に加熱し、焼鈍温度：Ａｃ１変態点未満で１．０ｈ以上保持する焼鈍を施す高炭素熱延鋼板の製造方法。     A method for manufacturing a high carbon hot rolled steel sheet according to any one of claims 1 to 4,
    Steel having the above-mentioned composition is subjected to hot rough rolling, finish rolling at a finish rolling end temperature: Ar 3 transformation point or higher, and then cooled to 650 to 750°C at an average cooling rate of 20 to 100°C/sec. ,
    Winding temperature: After winding at 500-700°C to make hot rolled steel sheet,
    A high-carbon hot-rolled steel sheet which is annealed by heating the hot-rolled steel sheet at an average heating rate of 15° C./h or more in a temperature range of 450 to 600° C. and holding it at an annealing temperature of less than Ac 1 transformation point for 1.0 hour or more. Manufacturing method.
  6. 　請求項１～４のいずれかに記載の高炭素熱延鋼板の製造方法であって、
    前記成分組成を有する鋼を、熱間粗圧延後、仕上圧延終了温度：Ａｒ３変態点以上で仕上圧延を行い、その後、平均冷却速度：２０～１００℃／ｓｅｃで６５０～７５０℃まで冷却し、
    巻取温度：５００～７００℃で巻き取り、熱延鋼板とした後、
    該熱延鋼板を、平均加熱速度：１５℃／ｈ以上で４５０～６００℃の温度範囲に加熱し、Ａｃ１変態点以上Ａｃ３変態点以下で０．５ｈ以上保持し、次いで平均冷却速度：１～２０℃／ｈでＡｒ１変態点未満に冷却し、Ａｒ１変態点未満で２０ｈ以上保持する焼鈍を施す高炭素熱延鋼板の製造方法。     A method for manufacturing a high carbon hot rolled steel sheet according to any one of claims 1 to 4,
    Steel having the above-mentioned composition is subjected to hot rough rolling, finish rolling at a finish rolling end temperature: Ar 3 transformation point or higher, and then cooled to 650 to 750°C at an average cooling rate of 20 to 100°C/sec. ,
    Winding temperature: After winding at 500-700°C to make hot rolled steel sheet,
    The hot-rolled steel sheet is heated to a temperature range of 450 to 600° C. at an average heating rate of 15° C./h or more and kept for 0.5 h or more at an Ac 1 transformation point or more and an Ac 3 transformation point or less, and then an average cooling rate: 1 ~ 20 ℃ / h with cooling to less than Ar 1 transformation point, the method of producing a high-carbon hot-rolled steel sheet subjected to annealing for holding 20h or less than Ar 1 transformation point.
PCT/JP2020/000783 2019-01-30 2020-01-14 High-carbon hot-rolled steel sheet and method for manufacturing same WO2020158357A1 (en)

Priority Applications (5)

Application Number Priority Date Filing Date Title
US17/425,824 US20220170126A1 (en) 2019-01-30 2020-01-14 High-carbon hot-rolled steel sheet and method for manufacturing the same
KR1020217023358A KR102570145B1 (en) 2019-01-30 2020-01-14 High-carbon hot-rolled steel sheet and manufacturing method thereof
EP20749360.2A EP3901303A4 (en) 2019-01-30 2020-01-14 High-carbon hot-rolled steel sheet and method for manufacturing same
CN202080011346.5A CN113366137B (en) 2019-01-30 2020-01-14 High carbon hot-rolled steel sheet and method for producing same
JP2020520327A JP6927427B2 (en) 2019-01-30 2020-01-14 High carbon hot-rolled steel sheet and its manufacturing method

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
JP2019-013957 2019-01-30
JP2019013957 2019-01-30

Publications (1)

Publication Number Publication Date
WO2020158357A1 true WO2020158357A1 (en) 2020-08-06

Family

ID=71842205

Family Applications (1)

Application Number Title Priority Date Filing Date
PCT/JP2020/000783 WO2020158357A1 (en) 2019-01-30 2020-01-14 High-carbon hot-rolled steel sheet and method for manufacturing same

Country Status (7)

Country Link
US (1) US20220170126A1 (en)
EP (1) EP3901303A4 (en)
JP (1) JP6927427B2 (en)
KR (1) KR102570145B1 (en)
CN (1) CN113366137B (en)
TW (1) TWI738186B (en)
WO (1) WO2020158357A1 (en)

Cited By (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP7444097B2 (en) 2021-02-10 2024-03-06 Ｊｆｅスチール株式会社 Hot rolled steel sheet and its manufacturing method

Families Citing this family (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP4317484A1 (en) * 2021-03-31 2024-02-07 JFE Steel Corporation Cold-rolled steel sheet and cold-rolled steel sheet manufacturing method
CN115181905B (en) * 2022-06-23 2023-09-15 首钢集团有限公司 Gear steel and production method thereof
CN115418567B (en) * 2022-08-31 2024-01-19 马鞍山钢铁股份有限公司 Nb-Ti-B microalloyed high-temperature-resistant low-internal oxidation carburized gear steel and manufacturing method thereof

Citations (12)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2004250767A (en) * 2003-02-21 2004-09-09 Nippon Steel Corp Steel for machine structural having excellent cold workability, and production method therefor
JP2005344194A (en) 2004-06-07 2005-12-15 Nippon Steel Corp High-carbon steel sheet having excellent workability and hardenability
JP4012475B2 (en) 2003-02-21 2007-11-21 新日本製鐵株式会社 Machine structural steel excellent in cold workability and low decarburization and method for producing the same
JP2009299189A (en) 2009-09-08 2009-12-24 Nisshin Steel Co Ltd High carbon steel sheet for precision blanking
WO2010106748A1 (en) * 2009-03-16 2010-09-23 新日本製鐵株式会社 Boron-containing steel sheet with excellent hardenability and method of manufacturing same
JP5458649B2 (en) 2009-04-28 2014-04-02 Ｊｆｅスチール株式会社 High carbon hot rolled steel sheet and manufacturing method thereof
JP2015017284A (en) 2013-07-09 2015-01-29 Ｊｆｅスチール株式会社 High carbon hot-rolled steel sheet excellent in hardenability and workability, and method for manufacturing the same
JP2015017283A (en) 2013-07-09 2015-01-29 Ｊｆｅスチール株式会社 High carbon hot-rolled steel sheet excellent in hardenability and workability, and method for manufacturing the same
WO2015146174A1 (en) * 2014-03-28 2015-10-01 Ｊｆｅスチール株式会社 High-carbon hot-rolled steel sheet and method for producing same
WO2015146173A1 (en) 2014-03-28 2015-10-01 Ｊｆｅスチール株式会社 High-carbon hot-rolled steel sheet and method for producing same
WO2016190396A1 (en) * 2015-05-26 2016-12-01 新日鐵住金株式会社 Steel sheet and method for producing same
WO2019131099A1 (en) * 2017-12-25 2019-07-04 Ｊｆｅスチール株式会社 Hot-rolled steel sheet and method for manufacturing same

Family Cites Families (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS5458649U (en) 1977-09-30 1979-04-23
JPS5788967A (en) 1980-11-21 1982-06-03 Showa Alum Corp Formation of porous layer on metallic surface
US20050199322A1 (en) * 2004-03-10 2005-09-15 Jfe Steel Corporation High carbon hot-rolled steel sheet and method for manufacturing the same
KR20140110994A (en) * 2012-01-05 2014-09-17 제이에프이 스틸 가부시키가이샤 High carbon hot-rolled steel sheet and method for producing same
CN105378133B (en) * 2013-07-09 2018-03-06 杰富意钢铁株式会社 High-carbon hot-rolled steel sheet and its manufacture method
WO2016190370A1 (en) * 2015-05-26 2016-12-01 新日鐵住金株式会社 Steel sheet and method for producing same
BR112017025756A2 (en) * 2015-06-17 2018-08-14 Nippon Steel & Sumitomo Metal Corporation A steel plate and a manufacturing method

Patent Citations (13)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP4012475B2 (en) 2003-02-21 2007-11-21 新日本製鐵株式会社 Machine structural steel excellent in cold workability and low decarburization and method for producing the same
JP2004250767A (en) * 2003-02-21 2004-09-09 Nippon Steel Corp Steel for machine structural having excellent cold workability, and production method therefor
JP2005344194A (en) 2004-06-07 2005-12-15 Nippon Steel Corp High-carbon steel sheet having excellent workability and hardenability
WO2010106748A1 (en) * 2009-03-16 2010-09-23 新日本製鐵株式会社 Boron-containing steel sheet with excellent hardenability and method of manufacturing same
JP4782243B2 (en) 2009-03-16 2011-09-28 新日本製鐵株式会社 Boron-added steel sheet with excellent hardenability and manufacturing method
JP5458649B2 (en) 2009-04-28 2014-04-02 Ｊｆｅスチール株式会社 High carbon hot rolled steel sheet and manufacturing method thereof
JP2009299189A (en) 2009-09-08 2009-12-24 Nisshin Steel Co Ltd High carbon steel sheet for precision blanking
JP2015017284A (en) 2013-07-09 2015-01-29 Ｊｆｅスチール株式会社 High carbon hot-rolled steel sheet excellent in hardenability and workability, and method for manufacturing the same
JP2015017283A (en) 2013-07-09 2015-01-29 Ｊｆｅスチール株式会社 High carbon hot-rolled steel sheet excellent in hardenability and workability, and method for manufacturing the same
WO2015146174A1 (en) * 2014-03-28 2015-10-01 Ｊｆｅスチール株式会社 High-carbon hot-rolled steel sheet and method for producing same
WO2015146173A1 (en) 2014-03-28 2015-10-01 Ｊｆｅスチール株式会社 High-carbon hot-rolled steel sheet and method for producing same
WO2016190396A1 (en) * 2015-05-26 2016-12-01 新日鐵住金株式会社 Steel sheet and method for producing same
WO2019131099A1 (en) * 2017-12-25 2019-07-04 Ｊｆｅスチール株式会社 Hot-rolled steel sheet and method for manufacturing same

Non-Patent Citations (1)

* Cited by examiner, † Cited by third party
Title
SATOSHI KINOSHIROTOMOHARU ISHIDAKUNIO INOSEKYOKO FUJIMOTO, TETSU-TO-HAGANE (IRON AND STEEL, vol. 99, no. 5, 2013, pages 362 - 365

Cited By (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP7444097B2 (en) 2021-02-10 2024-03-06 Ｊｆｅスチール株式会社 Hot rolled steel sheet and its manufacturing method

Also Published As

Publication number Publication date
CN113366137B (en) 2022-10-28
JP6927427B2 (en) 2021-08-25
US20220170126A1 (en) 2022-06-02
KR20210105417A (en) 2021-08-26
EP3901303A4 (en) 2021-11-03
TW202031912A (en) 2020-09-01
JPWO2020158357A1 (en) 2021-02-18
EP3901303A1 (en) 2021-10-27
CN113366137A (en) 2021-09-07
TWI738186B (en) 2021-09-01
KR102570145B1 (en) 2023-08-23

Similar Documents

Publication Publication Date Title
JP6927427B2 (en) High carbon hot-rolled steel sheet and its manufacturing method
CN108315637B (en) High carbon hot-rolled steel sheet and method for producing same
JP6244701B2 (en) High carbon hot rolled steel sheet excellent in hardenability and workability and method for producing the same
JP6065121B2 (en) High carbon hot rolled steel sheet and manufacturing method thereof
JP6569845B1 (en) High carbon hot rolled steel sheet and manufacturing method thereof
JPWO2015146173A1 (en) High carbon hot rolled steel sheet and manufacturing method thereof
JP6977880B2 (en) High carbon hot-rolled steel sheet and its manufacturing method
JP6402842B1 (en) High carbon hot rolled steel sheet and manufacturing method thereof
JP5884781B2 (en) High carbon hot rolled steel sheet excellent in hardenability and workability and method for producing the same
CN113490756B (en) Steel sheet, member, and method for producing same
JP2022122483A (en) Hot rolled steel sheet and method for producing the same
JP2022122482A (en) Hot rolled steel sheet and method for producing the same
CN118147535A (en) High-carbon hot-rolled steel sheet and method for producing same
JP2005272991A (en) Nitriding steel sheet having excellent wear resistance and press workability

Legal Events

Date Code Title Description
ENP Entry into the national phase

Ref document number: 2020520327

Country of ref document: JP

Kind code of ref document: A

121 Ep: the epo has been informed by wipo that ep was designated in this application

Ref document number: 20749360

Country of ref document: EP

Kind code of ref document: A1

ENP Entry into the national phase

Ref document number: 20217023358

Country of ref document: KR

Kind code of ref document: A

ENP Entry into the national phase

Ref document number: 2020749360

Country of ref document: EP

Effective date: 20210720

NENP Non-entry into the national phase

Ref country code: DE