WO2018186273A1 - Steel member, hot-rolled steel sheet for said steel member and production methods therefor - Google Patents

Steel member, hot-rolled steel sheet for said steel member and production methods therefor Download PDF

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Publication number
WO2018186273A1
WO2018186273A1 PCT/JP2018/013076 JP2018013076W WO2018186273A1 WO 2018186273 A1 WO2018186273 A1 WO 2018186273A1 JP 2018013076 W JP2018013076 W JP 2018013076W WO 2018186273 A1 WO2018186273 A1 WO 2018186273A1
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Prior art keywords
steel
hot
temperature
less
steel member
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PCT/JP2018/013076
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French (fr)
Japanese (ja)
Inventor
俊介 豊田
杉本　一郎
修司 川村
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Ｊｆｅスチール株式会社
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Priority to CA3057814A priority Critical patent/CA3057814C/en
Priority to KR1020197028866A priority patent/KR102319579B1/en
Priority to MX2019011941A priority patent/MX2019011941A/en
Priority to US16/500,613 priority patent/US20200190618A1/en
Priority to CN201880023595.9A priority patent/CN110494582B/en
Priority to JP2018536532A priority patent/JP6631715B2/en
Publication of WO2018186273A1 publication Critical patent/WO2018186273A1/en

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    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B1/00—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations
    • B21B1/22—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling plates, strips, bands or sheets of indefinite length
    • B21B1/24—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling plates, strips, bands or sheets of indefinite length in a continuous or semi-continuous process
    • B21B1/26—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling plates, strips, bands or sheets of indefinite length in a continuous or semi-continuous process by hot-rolling, e.g. Steckel hot mill
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21C—MANUFACTURE OF METAL SHEETS, WIRE, RODS, TUBES OR PROFILES, OTHERWISE THAN BY ROLLING; AUXILIARY OPERATIONS USED IN CONNECTION WITH METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL
    • B21C37/00—Manufacture of metal sheets, bars, wire, tubes or like semi-manufactured products, not otherwise provided for; Manufacture of tubes of special shape
    • B21C37/06—Manufacture of metal sheets, bars, wire, tubes or like semi-manufactured products, not otherwise provided for; Manufacture of tubes of special shape of tubes or metal hoses; Combined procedures for making tubes, e.g. for making multi-wall tubes
    • B21C37/08—Making tubes with welded or soldered seams
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—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
    • C21D8/10—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies
    • C21D8/105—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies 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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/08—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for tubular bodies or pipes
    • 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
    • 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
    • C21D9/50—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for welded joints
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/008—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
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—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
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • C—CHEMISTRY; METALLURGY
    • 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

Definitions

  • the present invention relates to a steel member, a hot-rolled steel sheet for the steel member, and a method for producing them. More specifically, the present invention relates to a steel member having excellent fatigue resistance in a plastic strain region, a hot-rolled steel sheet for the steel member, and a method for producing them.
  • the present invention relates to a welded steel pipe for coiled tubing, a welded steel pipe for line pipe, and a welded steel pipe for structural members for automobiles, which are required to have high strength and fatigue resistance in the plastic strain region, and in particular, welding for coiled tubing.
  • the present invention relates to a steel pipe and relates to an improvement in fatigue life in the plastic strain region of these steel members.
  • Patent Document 1 as a high-strength structural member and a driving force transmission member for an automobile or the like, or an oil-welded pipe for washing an oil well pipe, the yield strength after pipe forming is 700 MPa or more, the tensile strength is 800 MPa or more, and the elongation is 15% or more.
  • a method of manufacturing a high-strength electric resistance welded steel pipe having the following ductility is disclosed. According to this method, 0.09 to 0.18% C and a predetermined amount of Cu, Ni, Cr, and Mo alloy elements are contained, so that high-strength electric sewing that does not cause softening of the weld heat affected zone is achieved.
  • a steel pipe can be obtained.
  • a steel tube for coiled tubing that is required for fatigue use, in particular, fatigue resistance in the plastic strain region, has a problem of low durability life in repeated use.
  • Patent Document 2 discloses a steel strip for coiled tubing excellent in material uniformity and a method for manufacturing the same. According to this method, variation in yield strength in the coil width direction and the longitudinal direction is achieved by containing a predetermined amount of 0.10 to 0.16% C and an alloy element of Cr, Cu, Ni, Mo, Nb, and Ti. A steel strip for coiled tubing with a small diameter can be obtained. However, the fatigue resistance property in the plastic strain region is not sufficient, and there is a problem that the durability life in repeated use is low.
  • Patent Document 3 discloses a quenched and tempered steel pipe having excellent fatigue life for a steel pipe for a machine structure such as an automobile, particularly for a hollow stabilizer for an automobile. According to this method, a steel pipe having a high fatigue life can be obtained by containing a predetermined chemical component, setting the average particle size of the precipitated carbide to 0.5 ⁇ m or less, and setting the hardness at the center of the thickness to 400 HV.
  • the fatigue life level obtained with this steel pipe is a low stress-high cycle elastic region fatigue characteristic with a life of tens of thousands of cycles.
  • coiled tubing is used several hundred times while being repeatedly inserted and recovered into the well.
  • An object of the present invention is to provide a steel member having excellent fatigue resistance in a plastic strain region, a hot-rolled steel sheet as a raw material thereof, and a method for producing them.
  • the hot-rolled steel sheet used as the material of the steel member of the present invention is also referred to as “material hot-rolled steel sheet”.
  • the steel member of the present invention include steel pipes such as welded steel pipes and molded parts such as automobile structural members.
  • Examples of the welded steel pipe include a welded steel pipe for coiled tubing, a welded steel pipe for line pipes, and a welded steel pipe for structural members for automobiles.
  • the present inventors have made various changes in the chemical composition and production conditions of the hot-rolled steel sheet used as a material. Experiments were conducted. As a result, a steel having a specific chemical component is hot-rolled under a specific temperature processing condition or formed into a steel pipe shape and then heat-treated under a specific condition, thereby achieving high strength and an excellent plastic strain region. It has been found that a steel member that simultaneously satisfies fatigue resistance can be obtained.
  • the present invention has been completed based on such findings, and has the following configurations [1] to [9].
  • a steel member that contains 0.031 to 0.200% Ti by mass%, and 0.005% or more of Ti is precipitated as a precipitate having a particle size of 20 nm or less in the structure.
  • the steel member is, by mass%, C: 0.19 to 0.50%, Si: 0.002 to 1.5%, Mn: 0.4 to 2.5%, Al: 0.01 To 0.19%, Cr: 0.001 to 0.90%, B: 0.0001 to 0.0050%, Ti: 0.031 to 0.200%, P: 0.019% or less (0% S): 0.015% or less (including 0%), N: 0.008% or less (including 0%), O: 0.003% or less (including 0%), Sn: 0.10 % Or less (including 0%), the steel member according to [1] having a composition in which the balance is composed of Fe and inevitable impurities.
  • Nb 0.001 to 0.15%
  • V 0.001 to 0.15%
  • W 0.001 to 0.15%
  • Mo 0 by mass% 0.001 to 0.45%
  • Cu 0.001 to 0.45%
  • Ni 0.001 to 0.45%
  • Ca 0.0001 to 0.005%
  • Sb 0.0001 to 0.10
  • a method for producing a hot-rolled steel sheet for a steel member wherein the temperature range is cooled at an average cooling rate of 10 ° C./s or more and wound at a temperature of T Ti ⁇ 500 ° C. or less.
  • the steel member excellent in the fatigue resistance characteristic of a plastic strain area can be provided.
  • the hot-rolled steel sheet of the present invention is particularly suitable as a material for the steel member.
  • ADVANTAGE OF THE INVENTION According to this invention, the steel member which can make compatible the characteristic which is the strength and the fatigue-resistant characteristic in a plastic strain area at a high level can be provided. Therefore, as the steel member of the present invention, a welded steel pipe for coiled tubing, a welded steel pipe for line pipes, and a welded steel pipe for automotive structural members, which are particularly required to have high strength and fatigue resistance in the plastic strain region, are suitable. In particular, a welded steel pipe for coiled tubing is suitable.
  • the steel member of the present invention is obtained by subjecting a hot-rolled steel sheet (raw material hot-rolled steel sheet) produced by hot rolling under specific temperature processing conditions to a heat treatment under specific conditions.
  • a hot-rolled steel sheet raw material hot-rolled steel sheet
  • the heat treatment performed after forming the raw hot-rolled steel sheet is also referred to as “post-heat treatment”.
  • Ti 0.031 to 0.200% Ti precipitates as carbonitride in the hot rolling process, and suppresses recovery / recrystallization grain growth in the hot rolling process.
  • Ti By containing Ti, there is an effect that a desired fine ferrite phase particle size (1 to 50 ⁇ m) can be obtained in the structure (microstructure) of the hot-rolled steel sheet.
  • the refinement of the microstructure at the hot-rolled steel sheet stage leads to the refinement of the microstructure after heat treatment after subsequent forming (cold working) such as pipe forming and part forming, and excellent plasticity. Fatigue resistance in the strain range is obtained.
  • Tanaka et al. Proposed a model in which dislocations pile up irreversibly on the slip surface due to fatigue cycles, and when the stress generated at this time exceeds the critical stress, an initial crack occurs (reference: K. Tanaka and T. Mura: J Appl Mech., Vol. 48, p.97-103 (1981)).
  • G transverse elastic constant
  • Ws fracture energy per unit area
  • ⁇ Poisson's ratio
  • ⁇ decomposition shear stress range on the sliding surface
  • k dislocation frictional force on the sliding surface, etc.
  • the fatigue crack generation cycle Nc of each crystal grain becomes longer as the slip surface length d is shorter, that is, as the crystal grain size is smaller. Due to such a mechanism, it is considered that the refined microstructure material of the present invention is delayed in fatigue crack generation and exhibits excellent fatigue resistance characteristics in the plastic strain region.
  • Ti is strengthened by precipitation strengthening the matrix as a carbide, solid solution strengthening as a solid solution element, and strengthening transformation structure strengthening as a hardenability improving element, and after heat treatment after forming processing such as pipe making and part forming
  • This is an essential element that improves the strength of the steel and significantly improves the fatigue strength.
  • Such an effect is obtained when the Ti content is in the range of 0.031 to 0.200%, and when the Ti content is less than the lower limit of the above range, the effect is 0. 005% or more of Ti exists as solute Ti, and the heat treatment after the molding process cannot precipitate 0.005% or more of Ti as fine precipitates having a particle diameter of 20 nm or less, and the above effect is obtained.
  • the Ti content is in the range of 0.031 to 0.200%.
  • the Ti content is preferably more than 0.120%. Further, the Ti content is preferably 0.150% or less.
  • 0.005% or more of Ti is precipitated as a precipitate having a particle size of 20 nm or less.
  • the present inventors need a fatigue resistance property in a plastic strain region after a heat treatment (post heat treatment) performed after a forming process such as pipe making or part forming using a hot-rolled steel sheet as in the present invention. In that case, it is found that by applying post-heat treatment, 0.005% or more of Ti is precipitated as fine precipitates having a particle size of 20 nm or less, so that excellent fatigue resistance characteristics in the plastic strain region can be obtained. did.
  • C 0.19 to 0.50%
  • C is post-heat-treated under specific conditions to ensure high strength, and further binds to Ti during post-heat treatment, in particular, precipitates fine precipitates in the surface layer portion to cause fatigue resistance in the plastic strain region. It is an element that improves the characteristics.
  • the C content is less than 0.19%, it becomes difficult to obtain the desired strength (YS ⁇ 770 MPa) and fatigue resistance in the plastic strain region.
  • the C content exceeds 0.50%, the toughness and weldability of a steel member, for example, a steel pipe, cannot be ensured, so this is the upper limit. More preferably, the C content is more than 0.28%. More preferably, the C content is 0.30% or less.
  • Si 0.002 to 1.5%
  • Si is an element that improves fatigue resistance in the plastic strain region while ensuring a desired strength by solid solution strengthening. If the Si content is less than 0.002%, the strength is insufficient. On the other hand, if the content exceeds 1.5%, weldability deteriorates. Therefore, the Si content is preferably limited to 0.002 to 1.5%. More preferably, the Si content is 0.05% or more. More preferably, the Si content is 0.35% or less.
  • Mn 0.4 to 2.5%
  • Mn has a function of ensuring a desired strength by strengthening at low temperature during post-heat treatment and improving fatigue resistance in the plastic strain region. If the Mn content is less than 0.4%, this effect is not sufficiently exhibited. On the other hand, if the Mn content exceeds 2.5%, the weldability deteriorates. Therefore, the Mn content is preferably limited to 0.4 to 2.5%. More preferably, the Mn content is 1.09% or more. More preferably, the Mn content is 1.99% or less.
  • Al 0.01 to 0.19%
  • Al is a deoxidizing element during steel making, suppresses the growth of austenite grains in the hot rolling process, makes the crystal grains fine, and obtains a desired ferrite grain size (1 to 50 ⁇ m) after post-heat treatment, It has the function of improving fatigue resistance in the plastic strain region. If the Al content is less than 0.01%, these effects cannot be obtained, and the ferrite grain size becomes coarse. On the other hand, if the Al content exceeds 0.19%, the weldability deteriorates and the oxide type intervening The fatigue resistance tends to decrease due to the increase in the number of objects. More preferably, the Al content is 0.041% or more. More preferably, the Al content is 0.080% or less.
  • Cr 0.001 to 0.90% Cr has a function of securing a desired strength by strengthening at low temperature transformation during post-heat treatment and improving fatigue resistance in a plastic strain region. If the Cr content is less than 0.001%, this effect is not sufficiently exhibited. On the other hand, if the Cr content exceeds 0.90%, the weldability deteriorates. Therefore, the Cr content is preferably limited to 0.001 to 0.90%. More preferably, the Cr content is 0.001 to 0.19%.
  • B 0.0001 to 0.0050%
  • B has a function of ensuring a desired strength by strengthening at low temperature transformation during post-heat treatment and improving fatigue resistance in a plastic strain region. If the B content is less than 0.0001%, this effect is not sufficiently exhibited. On the other hand, if the B content exceeds 0.0050%, the fatigue resistance tends to decrease. Therefore, the B content is preferably limited to 0.0001 to 0.0050%. More preferably, the B content is 0.0005% or more. More preferably, the B content is 0.0035% or less.
  • P 0.019% or less (including 0%) P deteriorates fatigue resistance in the plastic strain region and deteriorates electroweldability through solidification co-segregation with Mn. If the P content exceeds 0.019%, the adverse effect becomes remarkable, so 0.019% is preferable as the upper limit.
  • S 0.015% or less (including 0%) S exists as an inclusion in steel as MnS or the like, and lowers fatigue resistance as a starting point of fatigue cracks in the plastic strain region.
  • the S content exceeds 0.015%, this adverse effect becomes significant, so it is preferable to set the upper limit at 0.015%. More preferably, the S content is 0.005% or less.
  • N 0.008% or less (including 0%) N forms Ti and TiN, precipitates as coarse precipitates, and consumes solid solution Ti.
  • N is added in the form of hot-rolled steel sheet at the raw material stage so that 0.005% or more of Ti is present as solute Ti, and heat treatment after forming is performed so that 0.005% or more of Ti is finely grained with a particle size of 20 nm or less. It precipitates as a good precipitate and reduces the effect of obtaining excellent fatigue resistance characteristics in the plastic strain region. If the N content exceeds 0.008%, this adverse effect becomes significant, so it is preferable to set the upper limit to 0.008%. More preferably, the N content is 0.0049% or less.
  • O 0.003% or less (including 0%) O exists as oxide inclusions and reduces the fatigue resistance of steel. If the content of O exceeds 0.003%, this adverse effect becomes remarkable, so 0.003% is preferably set as the upper limit. More preferably, the O content is 0.002% or less.
  • Sn 0.10% or less (including 0%) Sn exists as a solid solution element and reduces the hot ductility of steel. If the Sn content exceeds 0.10%, this adverse effect becomes significant, so it is preferable to set the upper limit to 0.10%. More preferably, the Sn content is 0.03% or less.
  • the balance is Fe and inevitable impurities.
  • the following elements can be further added for the purpose of improving the effects of the present invention.
  • Nb 0.001 to 0.15% Nb precipitates as a carbide, suppresses recovery / recrystallization grain growth in the hot rolling process, and has the effect of obtaining a desired ferrite grain size (1 to 50 ⁇ m), and can be contained as needed. If the Nb content is less than 0.001%, these effects cannot be obtained. On the other hand, when the content of Nb exceeds 0.15%, coarse precipitates are deposited on the surface layer portion due to strain-induced precipitation during hot rolling, and fine precipitates on the surface layer portion are reduced. Since fatigue resistance is reduced, the upper limit is made 0.15%. Therefore, when Nb is contained, the Nb content is set to 0.001 to 0.15%. More preferably, the Nb content is 0.001 to 0.009%.
  • V 0.001 to 0.15%
  • V precipitates as a carbide, suppresses recovery / recrystallization grain growth in the hot rolling process, and has the effect of obtaining a desired ferrite grain size (1 to 50 ⁇ m), and can be contained as necessary. If the V content is less than 0.001%, these effects cannot be obtained.
  • the content of V exceeds 0.15%, coarse precipitates are deposited on the surface layer portion due to strain-induced precipitation during hot rolling, and fine precipitates on the surface layer portion are reduced, and in the plastic strain region.
  • the upper limit is 0.15% because the fatigue resistance is reduced. Therefore, when V is contained, the content of V is set to 0.001 to 0.15%. More preferably, the V content is 0.001 to 0.049%.
  • W 0.001 to 0.15% W precipitates as carbide, suppresses recovery / recrystallization grain growth in the hot rolling process, supplements the effect of obtaining the desired ferrite grain size (1 to 50 ⁇ m), and is contained if necessary it can. If the W content is less than 0.001%, these effects cannot be obtained. On the other hand, when the content of W exceeds 0.15%, coarse precipitates are deposited on the surface layer portion due to strain-induced precipitation during hot rolling, and fine precipitates on the surface layer portion are reduced. The upper limit is 0.15% because the fatigue resistance is reduced. Therefore, when W is contained, the W content is set to 0.001 to 0.15%. More preferably, the W content is 0.001 to 0.049%.
  • Mo 0.001 to 0.45%
  • Mo has a function of securing a desired strength by low-temperature transformation strengthening or precipitation strengthening during post-heat treatment and complementing the effect of improving fatigue resistance in the plastic strain region, and can be contained as necessary. If the Mo content is less than 0.001%, this effect does not appear. On the other hand, if the Mo content exceeds 0.45%, the weldability deteriorates. Therefore, when Mo is contained, the Mo content is set to 0.001 to 0.45%. More preferably, the Mo content is 0.001 to 0.30%.
  • Cu and Ni are elements that have a function of complementing the effect of improving the fatigue strength of Mn, and at the same time, have the effect of increasing the corrosion resistance of the steel material, and can contain Cu and Ni as needed. These effects are manifested when the content of Cu and Ni is 0.001% or more. However, when the content exceeds 0.45% for Cu and Ni, the upper limit is 0.45% for lowering weldability. Therefore, when Cu is contained, the Cu content is set to 0.001 to 0.45%. When Ni is contained, the Ni content is set to 0.001 to 0.45%. More preferably, any element is 0.35% or less.
  • Ca 0.0001 to 0.005%
  • Ca has a so-called form control effect in which expanded MnS is granular Ca (Al) S (O), has the effect of suppressing fatigue cracking and improving fatigue resistance, and can be contained if necessary. .
  • This effect is manifested with a content of 0.0001% or more.
  • the content exceeding 0.005% is limited to 0.005% because the fatigue resistance is lowered by the increase of nonmetallic inclusions. Therefore, when Ca is contained, the content of Ca is set to 0.0001 to 0.005%.
  • Sb 0.0001 to 0.10%
  • Sb preferentially segregates on the surface, suppresses the intrusion of N from the atmosphere in the hot rolling process or the post heat treatment process, and functions to suppress a decrease in the effect of adding B due to the formation of BN.
  • This effect is manifested at a content of 0.0001% or more, but even if it exceeds 0.10%, the effect is saturated, so 0.10% is made the upper limit. Therefore, when Sb is contained, the Sb content is set to 0.0001 to 0.10%. More preferably, the Sb content is 0.0001 to 0.030%.
  • the average crystal grain size of the ferrite phase from the surface after post heat treatment to the plate thickness direction of 200 ⁇ m is 1 to 50 ⁇ m, and the particle size of 1 in the ferrite phase from the surface to the plate thickness direction of 200 ⁇ m.
  • the difference between the average hardness from the surface to the thickness direction of 200 ⁇ m and the average hardness in the vicinity of the thickness center excluding the central segregation part (absolute value) has a structure in which Ti carbide of 0.0 to 20 nm is precipitated.
  • the hardness (HV) is desirably ⁇ HV 50 points or less.
  • the microstructure of the steel member, the precipitation state of the precipitates, and the cross-sectional hardness are important for ensuring fatigue resistance in an excellent plastic strain region. If the average crystal grain size of the ferrite phase from the surface after post-heat treatment to the plate thickness direction of 200 ⁇ m exceeds 50 ⁇ m, the initial fatigue cracks are early and large, making it difficult to ensure fatigue resistance characteristics in a desired plastic strain region. On the other hand, since it is difficult industrially and economically to make the average crystal grain size of the ferrite phase less than 1 ⁇ m after the post heat treatment, this is set as the lower limit.
  • ferrite phase refers to the body phase iron of a body-centered cubic lattice, so-called polygonal ferrite, acicular ferrite, Widmanstatten ferrite, bainitic ferrite, bainite, and low carbon (C content of 1% or less). Includes a martensite organization.
  • the second phase other than the ferrite phase include austenite, carbide, pearlite, and high carbon martensite (C content exceeding 1%).
  • the structure of the steel member of the present invention preferably has the ferrite phase as a main phase.
  • the main phase refers to a phase occupying 51% or more by volume ratio, preferably 80% or more, and may be 100%.
  • the Ti carbide dimension in the ferrite phase from the surface to the plate thickness direction of 200 ⁇ m is important for ensuring the surface hardness and the fatigue resistance in a high plastic strain region.
  • the precipitation of 1.0 to 20 nm Ti carbide in the ferrite phase from the surface to the thickness direction of 200 ⁇ m suppresses the occurrence of initial fatigue cracks, reduces the size, and provides excellent resistance to plastic strain. The fatigue characteristics can be further improved.
  • the precipitation amount of Ti carbide of 1.0 to 20 nm is not particularly defined here. In addition to Ti carbide having a thickness of 1.0 to 20 nm, it is allowed to deposit Ti carbides having different dimensions.
  • the difference between the average hardness from the surface to the thickness direction of 200 ⁇ m and the average hardness in the vicinity of the thickness center excluding the center segregation part is ⁇ HV50 points or less, in order to ensure excellent fatigue resistance in the plastic strain region. is important.
  • the difference between the average hardness from the surface to the thickness direction of 200 ⁇ m and the average hardness in the vicinity of the center of the plate thickness excluding the center segregation part exceeds ⁇ HV50 points, the initial fatigue cracks occur quickly and greatly, and in the desired plastic strain region It is difficult to ensure the fatigue resistance characteristics. For this reason, it is desirable that the difference between the average hardness from the surface to the plate thickness direction of 200 ⁇ m and the average hardness in the vicinity of the plate thickness center excluding the center segregation portion is ⁇ HV 50 points or less.
  • the difference between the average hardness from the surface to the thickness direction of 200 ⁇ m and the average hardness in the vicinity of the thickness center excluding the center segregation part is the micro Vickers hardness at a pitch of 25 ⁇ m in the thickness direction between 50 and 200 ⁇ m in the thickness direction.
  • HV (0.1) averaged 7 points HV (0.1) S , avoiding the center segregation around the center of the plate thickness
  • HV (0.1 average value HV (0.1) the difference in C) was measured 7 points was measured as HV (0.1) C -HV (0.1 ) S.
  • the hot-rolled steel sheet (raw material hot-rolled steel sheet) for steel members of the present invention is particularly suitable for obtaining the steel member of the present invention.
  • the material hot-rolled steel sheet of the present invention contains 0.031 to 0.200% Ti by mass%, and 0.005% or more of Ti exists in the structure as solute Ti. Thereby, after performing a predetermined heat treatment after forming, 0.005% or more of Ti can be precipitated as fine precipitates having a particle size of 20 nm or less in the structure of the steel member, and in the plastic strain region It is possible to obtain a steel member that is excellent in fatigue resistance and also excellent in strength characteristics.
  • the composition of the material hot-rolled steel sheet of the present invention is the same as the composition of the steel member.
  • the thickness of the tip and tail ends, which are both ends in the longitudinal direction is 5 compared to the thickness of the intermediate portion (longitudinal central portion) other than both ends in the longitudinal direction.
  • it is ⁇ 50% thick.
  • the temperature is the surface temperature of a steel slab or the like.
  • a steel slab obtained by casting steel having the above composition is used as a starting material.
  • the production method of the starting material is not particularly limited.
  • the molten steel having the above composition is melted by a conventional melting method such as a converter, and a steel slab is obtained by a normal casting method such as a continuous casting method. Is mentioned.
  • the material hot-rolled steel sheet of the present invention can be manufactured by hot rolling a steel slab containing 0.031 to 0.200% Ti under predetermined conditions.
  • the amount of dissolved Ti is less than 0.005%, and the fatigue resistance characteristics in the plastic strain region that are remarkably excellent after post-heat treatment cannot be obtained.
  • T Ti equilibrium solid solution temperature
  • the slab extraction temperature is preferably 1620 K or less from the viewpoint of preventing the crystal grain size from becoming coarse, and the slab leveling is ensured from the viewpoint of ensuring uniformity of the solid solution state of Ti and sufficient solid solution time.
  • the heat time time for holding the slab at a temperature higher than the equilibrium solid solution temperature T Ti ) is preferably 10 min or more.
  • T Ti -400 ° C or higher finish rolling temperature When hot rolling finish rolling temperature is lower than T Ti -400 ° C, strain induced precipitation due to additional shear strain by upper and lower rolls near the surface or heat removal by rolls and cooling water.
  • the amount of solid solution Ti present in the vicinity of the front surface is less than 0.005% at the stage of the raw hot-rolled steel sheet, and fatigue resistance in a plastic strain region that is remarkably excellent after post-heat treatment Characteristics are not obtained.
  • 0.005% or more of Ti including the vicinity of the surface is present as solute Ti at the stage of the raw hot rolled steel sheet, and it is reduced to 0 by heat treatment after forming.
  • 0.005% or more of Ti can be precipitated as fine precipitates having a particle size of 20 nm or less, and a particularly excellent fatigue resistance property in a plastic strain region can be obtained.
  • Ti including the vicinity of the surface is solidified at the stage of the hot rolled steel sheet. It exists as molten Ti, and 0.005% or more of Ti can be precipitated as fine precipitates with a particle size of 20 nm or less by heat treatment after forming, and has excellent fatigue resistance characteristics in the plastic strain region. can get.
  • Winding temperature of T Ti ⁇ 500 ° C. or lower When the winding temperature exceeds T Ti ⁇ 500 ° C., precipitation of Ti precipitates is promoted before coil cooling, and solid solution Ti present at the stage of the raw hot rolled steel sheet The amount is less than 0.005%, and the fatigue resistance property in the plastic strain region which is remarkably excellent after post heat treatment cannot be obtained.
  • T Ti ⁇ 500 ° C. or less By setting the winding temperature to T Ti ⁇ 500 ° C. or less, 0.005% or more of Ti including the vicinity of the surface is present as solute Ti at the stage of the raw hot-rolled steel sheet.
  • the finish rolling temperature and the coiling temperature are surface temperatures at the center of the coil width, and the average cooling rate is obtained from the surface temperature.
  • a hot-rolled steel sheet (raw material hot-rolled steel sheet) in which 0.005% or more of Ti exists as a solid solution Ti in the structure is obtained.
  • the steel member of the present invention is manufactured by subjecting the material hot-rolled steel sheet to a predetermined heat treatment after forming.
  • a predetermined heat treatment for example, if a steel member is a steel pipe, a pipe making process will be mentioned. If the steel member is a welded steel pipe, the welding process may be performed after the pipe making process.
  • the steel member is a molded part such as a structural member for an automobile, press working or the like can be given. After forming, heat treatment is performed under the following conditions.
  • the solid solution Ti is not precipitated as fine precipitates having a thickness of 20 nm or less, and the remarkably excellent fatigue resistance property in the plastic strain region cannot be obtained.
  • the heating temperature exceeds 1050 ° C.
  • the particle size of the ferrite phase exceeds 50 ⁇ m, and it becomes difficult to obtain fatigue resistance characteristics in a plastic strain region that is remarkably excellent. If the cooling rate in the temperature range of 550 to 400 ° C. is less than 10 ° C./s, sufficient strength (YS ⁇ 770 MPa) cannot be obtained.
  • the heating temperature is more preferably in the range of 700 to 1000 ° C.
  • the raw hot-rolled steel sheet is left as it is, or, if necessary, pickling, cold rolling, annealing, plating, or a plurality of After processing, slitting to a predetermined plate width, one or more coils are welded and joined in the longitudinal direction, formed into a roughly circular cross-section by roll forming or press forming, and the end is subjected to high-frequency electric seam welding, laser Joined by welding, etc., and heated online or offline to a temperature exceeding 550 ° C. and below 1050 ° C., cooling the temperature range of 550 to 400 ° C. at an average cooling rate of 10 ° C./s or more to form a coil.
  • the raw hot-rolled steel sheet is left as it is, or after performing any one or more of pickling, cold rolling, annealing, plating as necessary.
  • 0.005% or more of Ti precipitates as fine precipitates having a particle size of 20 nm or less, and the fatigue resistance characteristics in the plastic strain region that are remarkably excellent are obtained.
  • Example 1 A steel slab having the composition shown in Table 1 (steel types C to L) was extracted from a heating furnace at a slab surface temperature of about 1220 ° C and a slab center temperature of about 1210 ° C, and finish rolling reduction: 91%, coil width center finish rolling temperature of about 860 ° C., coil width direction minimum finish rolling temperature of about 850 ° C., and cooled at an average cooling rate of T Ti -400 ° C. from T Ti -500 temperature range of about 20 ° C.
  • sheet thickness: about 5 mm, the thickness of the front and rear end portions is about 10% thicker than the longitudinal center portion was obtained by hot rolling at a temperature of 0 ° C. (No. 3 to 12).
  • a raw hot-rolled steel sheet was obtained in the same manner as above ( No.
  • a welded steel pipe having an outer diameter of 50.8 mm and a thickness of about 5 mm was obtained.
  • the whole welded steel pipe is continuously heated at high frequency, heated at a heating temperature of 920 ° C and a holding time of about 5 seconds, then cooled with water from the outer surface, and the temperature range of 550 to 400 ° C is an average of about 50 ° C / s.
  • a heat treatment for cooling at a cooling rate was performed.
  • Specimens were collected from these welded steel pipes and subjected to a structure observation test, a precipitate, a quantitative test of the solid solution amount, a tensile test, a plastic strain region fatigue test, and a low temperature toughness test.
  • the test method was as follows.
  • Microstructure observation test Samples of microstructural observation specimens were collected so that the circumferential cross section of these welded steel pipes became the observation surface, polished, nital-corroded, and observed with a scanning electron microscope (3000 times).
  • the average grain size of the ferrite phase was determined by an EBSD (Electron BackScatter Diffraction) method with an inclination angle of 15 ° or more with adjacent grains as a grain boundary.
  • the average particle diameter from the surface to the plate thickness direction of 200 ⁇ m was measured by averaging three points at a pitch of 50 ⁇ m between the plate thickness directions of 50 to 200 ⁇ m and the center segregation portion around the plate thickness center, Values obtained by measuring and averaging three points at a pitch of 50 ⁇ m in the plate thickness direction were obtained.
  • the precipitate was peeled from the sample piece and extracted into an aqueous SHMP solution.
  • the SHMP aqueous solution containing the precipitate is filtered using a filter in the order of hole diameters of 100 nm and 20 nm, and the residue and filtrate on the filtered filter are analyzed using an ICP emission spectroscopic analyzer.
  • the absolute amount of Ti in the residue and the filtrate is measured, and the absolute amount of Ti contained in precipitates having a particle size of more than 100 nm, precipitates having a particle size of 100 nm or less and more than 20 nm, and precipitates having a particle size of 20 nm or less. , Tisp was obtained respectively.
  • the electrolytic mass was calculated
  • the electrolytic solution after electrolysis was used as an analysis solution, and the concentration of Ti and Ti as a comparative element in the solution was measured using ICP mass spectrometry. Based on the obtained concentration, the concentration ratio of Ti to Fe was calculated, and the content ratio of Ti in a solid solution state was determined by multiplying the content ratio of Fe in the sample. In addition, the content rate of Fe in a sample can be calculated
  • the quantitative test of the precipitate and the solid solution amount was also performed on the welded steel pipe before the post-heat treatment.
  • the average hardness (HV (0.1) S ) from the surface to the thickness direction of 200 ⁇ m and the average hardness near the center of the plate thickness (HV (0.1) C ) excluding the central segregation part are measured by the above method.
  • a difference ⁇ HV (HV (0.1) C ⁇ HV (0.1) S ) between the average hardness in the thickness direction of 200 ⁇ m and the average hardness in the vicinity of the thickness center excluding the center segregation portion was determined. The obtained results are shown in Table 2.
  • the number of cycles in the above-described plastic strain region fatigue test is 1000 cycles or more, and the fatigue resistance property in the plastic strain region is excellent.
  • YS is 770 MPa or more and excellent in strength characteristics.
  • the Charpy fracture surface transition temperature is ⁇ 30 ° C. or lower, and the low temperature toughness is excellent.
  • the composition of steel does not satisfy the scope of the present invention, and Ti deposited as a precipitate having a particle size of 20 nm or less is less than 0.005%.
  • steel component composition does not meet the scope of the present invention.
  • No. 13 does not provide fatigue resistance characteristics in the desired plastic strain region.
  • Example 2 A steel slab having the composition of steel types A, B, and C shown in Table 1 is subjected to hot rolling under the conditions shown in Table 3 and a hot-rolled steel sheet (thickness: about 5 mm, the thickness at the front and rear ends is the longitudinal center) About 10% thicker).
  • a hot-rolled steel sheet thickness: about 5 mm, the thickness at the front and rear ends is the longitudinal center
  • the open pipe is electro-welded by high-frequency resistance welding, and the width drawing ratio is 4%.
  • a welded steel pipe having an outer diameter of 50.8 mm and a thickness of about 5 mm was obtained.
  • the entire welded steel pipe was continuously heated at a high frequency and heat treated under the conditions shown in Table 3.
  • Specimens were collected from these welded steel pipes, and subjected to a structure observation test, a precipitate, a quantitative test of the solid solution amount, a tensile test, a plastic strain region fatigue test, a low temperature toughness test, and a Vickers hardness measurement.
  • No. for No. 23 the hot-rolled steel sheet was pickled, blanked to a predetermined size, and subjected to a heat treatment under the conditions shown in Table 3 for a molded part by pressing. And the test piece was extract
  • Table 4 shows the obtained results. In Tables 3 and 4, the above No. Results 1 to 3 are also shown.
  • the number of cycles in the plastic strain region fatigue test is 1000 cycles or more, and the fatigue resistance property in the plastic strain region is excellent.
  • YS is 770 MPa or more and excellent in strength characteristics.
  • the Charpy fracture surface transition temperature is ⁇ 30 ° C. or lower, and the low temperature toughness is excellent.
  • the amount of Ti deposited as a precipitate having a particle size of 20 nm or less is outside the scope of the present invention. In Nos. 14 to 20, fatigue resistance characteristics in a desired plastic strain region cannot be obtained.

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Abstract

The purpose of the present invention is to provide a steel member with excellent fatigue resistance characteristics in the plastic strain range, a hot-rolled steel sheet as the material therefor and production methods therefor. A steel member, which contains 0.031-0.200 mass% Ti and in the structure of which at least 0.005 mass% of the Ti is precipitated as precipitates with a particle size of 20 nm or less. A hot-rolled steel sheet for said steel member that contains 0.031-0.200 mass% Ti and in the structure of which at least 0.005 mass% of the Ti is present as Ti solid solution. A production method for said steel member in which, after molding of the hot-rolled steel sheet, heat treatment is performed wherein the member is heated to a temperature greater than 550°C and not exceeding 1050°C and then cooled at an average cooling rate of at least 10°C/s for the temperature range of 550-400°C. A production method for said hot-rolled steel sheet in which a steel slab containing 0.031-0200 mass% Ti is extracted under temperature conditions higher than the equilibrium solid solution temperature TTi determined from a specified formula, after which finish-rolling is completed at a temperature of at least TTi-400°C, the sheet is cooled at an average cooling rate of at least 10°C/s for the temperature range from TTi-400°C to TTi-500°C, and is coiled at a temperature of TTi-500°C or less.

Description

鋼部材、前記鋼部材用の熱延鋼板およびこれらの製造方法Steel member, hot-rolled steel sheet for the steel member, and production method thereof
　本発明は、鋼部材、前記鋼部材用の熱延鋼板およびこれらの製造方法に関する。本発明は、より具体的には、塑性歪域の耐疲労特性に優れた鋼部材、前記鋼部材用の熱延鋼板およびこれらの製造方法に関する。本発明は、特に、高強度で塑性歪域の耐疲労特性を要求される、コイルドチュービング用溶接鋼管、ラインパイプ用溶接鋼管、自動車用構造部材用溶接鋼管に関し、なかでもコイルドチュービング用溶接鋼管に係るものであり、これらの鋼部材の塑性歪域での疲労寿命の改善に関するものである。 The present invention relates to a steel member, a hot-rolled steel sheet for the steel member, and a method for producing them. More specifically, the present invention relates to a steel member having excellent fatigue resistance in a plastic strain region, a hot-rolled steel sheet for the steel member, and a method for producing them. In particular, the present invention relates to a welded steel pipe for coiled tubing, a welded steel pipe for line pipe, and a welded steel pipe for structural members for automobiles, which are required to have high strength and fatigue resistance in the plastic strain region, and in particular, welding for coiled tubing. The present invention relates to a steel pipe and relates to an improvement in fatigue life in the plastic strain region of these steel members.
　特許文献１には、自動車等の高強度構造部材および駆動力伝達部材、あるいは油井管洗浄用電縫管として、造管後の降伏強度７００ＭＰａ以上、引張強度８００ＭＰａ以上の強度と、伸び１５％以上の延性を有する高張力電縫鋼管の製造方法が開示されている。この方法によれば、０．０９～０．１８％のＣと、Ｃｕ、Ｎｉ、Ｃｒ、Ｍｏの合金元素を所定量含有することで、溶接熱影響部の軟化をもたらすことない高張力電縫鋼管を得ることができる。しかしながら、疲労用途、特に塑性歪域の耐疲労特性を要求されるコイルドチュービング用鋼管としては、繰返し使用での耐久寿命が低いという問題があった。 In Patent Document 1, as a high-strength structural member and a driving force transmission member for an automobile or the like, or an oil-welded pipe for washing an oil well pipe, the yield strength after pipe forming is 700 MPa or more, the tensile strength is 800 MPa or more, and the elongation is 15% or more. A method of manufacturing a high-strength electric resistance welded steel pipe having the following ductility is disclosed. According to this method, 0.09 to 0.18% C and a predetermined amount of Cu, Ni, Cr, and Mo alloy elements are contained, so that high-strength electric sewing that does not cause softening of the weld heat affected zone is achieved. A steel pipe can be obtained. However, a steel tube for coiled tubing that is required for fatigue use, in particular, fatigue resistance in the plastic strain region, has a problem of low durability life in repeated use.
　特許文献２には、材質均一性に優れたコイルドチュービング用鋼帯およびその製造方法が開示されている。この方法によれば、０．１０～０．１６％のＣと、Ｃｒ、Ｃｕ、Ｎｉ、Ｍｏ、Ｎｂ、Ｔｉの合金元素を所定量含有することでコイル幅方向、長手方向の降伏強度のばらつきが小さいコイルドチュービング用鋼帯を得ることができる。しかしながら、塑性歪域の耐疲労特性は十分ではなく、繰返し使用での耐久寿命が低いという問題があった。 Patent Document 2 discloses a steel strip for coiled tubing excellent in material uniformity and a method for manufacturing the same. According to this method, variation in yield strength in the coil width direction and the longitudinal direction is achieved by containing a predetermined amount of 0.10 to 0.16% C and an alloy element of Cr, Cu, Ni, Mo, Nb, and Ti. A steel strip for coiled tubing with a small diameter can be obtained. However, the fatigue resistance property in the plastic strain region is not sufficient, and there is a problem that the durability life in repeated use is low.
　特許文献３には、自動車等の機械構造物用鋼管、特に自動車用中空スタビライザー用として、疲労寿命の優れた焼入れ・焼戻し鋼管が開示されている。この方法によれば、所定の化学成分を含有し、析出炭化物の平均粒径を０．５μｍ以下とし、肉厚中心部の硬さを４００ＨＶとすることにより、高疲労寿命の鋼管が得られる。しかしながら、本鋼管で得られる疲労寿命レベルは、寿命が数万サイクルとなる低応力－高サイクルの弾性域疲労特性である。一方、コイルドチュービングは、抗井への挿入、回収を繰り返しながら数百回使用される。コイルの巻き戻し－巻き付けならびに、抗井へ挿入する際の湾曲（グースネック）部分では２％程度の塑性域の歪が加わり、１００～１０００サイクルの高歪－低サイクル疲労強度が必要となる。一般に、弾性域疲労のように応力振幅一定の条件での疲労強度は、材料強度を上げることで増加する。一方、コイルドチュービングに加わる長手方向歪は、コイルとグースネックの内径で決まる歪一定条件に相当し、所謂Morrowの式の疲労延性係数の寄与が大きくなるため、高強度化は必ずしも寿命の向上につながらず所望の塑性歪域の耐疲労特性が得られないという問題があった。 Patent Document 3 discloses a quenched and tempered steel pipe having excellent fatigue life for a steel pipe for a machine structure such as an automobile, particularly for a hollow stabilizer for an automobile. According to this method, a steel pipe having a high fatigue life can be obtained by containing a predetermined chemical component, setting the average particle size of the precipitated carbide to 0.5 μm or less, and setting the hardness at the center of the thickness to 400 HV. However, the fatigue life level obtained with this steel pipe is a low stress-high cycle elastic region fatigue characteristic with a life of tens of thousands of cycles. On the other hand, coiled tubing is used several hundred times while being repeatedly inserted and recovered into the well. When the coil is unwound and wound, and the curved portion (gooseneck) is inserted into the well, a strain in the plastic region of about 2% is applied, and high strain and low cycle fatigue strength of 100 to 1000 cycles are required. In general, the fatigue strength under the condition of a constant stress amplitude such as elastic region fatigue increases by increasing the material strength. On the other hand, the longitudinal strain applied to the coiled tubing corresponds to a constant strain condition determined by the inner diameter of the coil and the gooseneck, and the contribution of the fatigue ductility factor of the so-called Morrow equation becomes large. There is a problem that fatigue resistance characteristics in a desired plastic strain region cannot be obtained.
特許第３４９１３３９号公報Japanese Patent No. 3491339 特許第５４９４８９５号公報Japanese Patent No. 5494895 特許第５１９６９３４号公報Japanese Patent No. 5196934
　本発明は、塑性歪域での耐疲労特性に優れた鋼部材と、その素材となる熱延鋼板およびこれらの製造方法を提供することを目的とする。 An object of the present invention is to provide a steel member having excellent fatigue resistance in a plastic strain region, a hot-rolled steel sheet as a raw material thereof, and a method for producing them.
　なお、本発明でいう、「塑性歪域の耐疲労特性に優れた」ないし「優れた塑性歪域の耐疲労特性」とは、引張モード、ひずみ制御モード、ひずみ比＝０、全ひずみ範囲２．０％の条件で引張疲労試験を行った場合の破断までの繰り返し数が１０００回以上である場合をいうものとする。
　また、本発明の鋼部材の素材となる熱延鋼板を、「素材熱延鋼板」ともいう。
　本発明の鋼部材としては、溶接鋼管等の鋼管、自動車用構造部材等の成形部品等が挙げられる。溶接鋼管としては、コイルドチュービング用溶接鋼管、ラインパイプ用溶接鋼管、自動車用構造部材用溶接鋼管等が挙げられる。 In the present invention, “excellent fatigue resistance property in the plastic strain region” or “excellent fatigue resistance property in the plastic strain region” means tensile mode, strain control mode, strain ratio = 0, total strain range 2 When the tensile fatigue test is performed under the condition of 0.0%, the number of repetitions until breakage is 1000 times or more.
Moreover, the hot-rolled steel sheet used as the material of the steel member of the present invention is also referred to as “material hot-rolled steel sheet”.
Examples of the steel member of the present invention include steel pipes such as welded steel pipes and molded parts such as automobile structural members. Examples of the welded steel pipe include a welded steel pipe for coiled tubing, a welded steel pipe for line pipes, and a welded steel pipe for structural members for automobiles.
　本発明者らは、強度と塑性歪域での耐疲労特性という、相反する特性を高度なレベルで両立させるために、素材となる熱延鋼板の化学成分、製造条件を種々変化させて系統的な実験検討を行なった。その結果、特定化学成分を有する鋼を、特定温度加工条件で熱間圧延し、或いは鋼管形状等に成形加工したのちに特定の条件で熱処理することで、高い強度と優れた塑性歪域での耐疲労特性を同時に満たす鋼部材が得られることを見出した。
　本発明はこのような知見に基づいて完成されたものであり、以下の［１］～［９］の構成を有する。 In order to make the conflicting properties of strength and fatigue resistance in the plastic strain range compatible at a high level, the present inventors have made various changes in the chemical composition and production conditions of the hot-rolled steel sheet used as a material. Experiments were conducted. As a result, a steel having a specific chemical component is hot-rolled under a specific temperature processing condition or formed into a steel pipe shape and then heat-treated under a specific condition, thereby achieving high strength and an excellent plastic strain region. It has been found that a steel member that simultaneously satisfies fatigue resistance can be obtained.
The present invention has been completed based on such findings, and has the following configurations [1] to [9].
［１］質量％で、Ｔｉを０．０３１～０．２００％含有し、組織中に０．００５％以上のＴｉが粒径２０ｎｍ以下の析出物として析出している、鋼部材。
［２］前記鋼部材は、質量％で、Ｃ：０．１９～０．５０％、Ｓｉ：０．００２～１．５％、Ｍｎ：０．４～２．５％、Ａｌ：０．０１～０．１９％、Ｃｒ：０．００１～０．９０％、Ｂ：０．０００１～０．００５０％、Ｔｉ：０．０３１～０．２００％、Ｐ：０．０１９％以下（０％を含む）、Ｓ：０．０１５％以下（０％を含む）、Ｎ：０．００８％以下（０％を含む）、Ｏ：０．００３％以下（０％を含む）、Ｓｎ：０．１０％以下（０％を含む）を含有し、残部がＦｅおよび不可避的不純物からなる組成を有する、［１］に記載の鋼部材。
［３］前記組成に加えてさらに、質量％で、Ｎｂ：０．００１～０．１５％、Ｖ：０．００１～０．１５％、Ｗ：０．００１～０．１５％、Ｍｏ：０．００１～０．４５％、Ｃｕ：０．００１～０．４５％、Ｎｉ：０．００１～０．４５％、Ｃａ：０．０００１～０．００５％、Ｓｂ：０．０００１～０．１０％のうちから選ばれた１種または２種以上を含有する、［２］に記載の鋼部材。
［４］前記鋼部材が溶接鋼管である、［１］～［３］のいずれかに記載の鋼部材。 [1] A steel member that contains 0.031 to 0.200% Ti by mass%, and 0.005% or more of Ti is precipitated as a precipitate having a particle size of 20 nm or less in the structure.
[2] The steel member is, by mass%, C: 0.19 to 0.50%, Si: 0.002 to 1.5%, Mn: 0.4 to 2.5%, Al: 0.01 To 0.19%, Cr: 0.001 to 0.90%, B: 0.0001 to 0.0050%, Ti: 0.031 to 0.200%, P: 0.019% or less (0% S): 0.015% or less (including 0%), N: 0.008% or less (including 0%), O: 0.003% or less (including 0%), Sn: 0.10 % Or less (including 0%), the steel member according to [1] having a composition in which the balance is composed of Fe and inevitable impurities.
[3] In addition to the above composition, Nb: 0.001 to 0.15%, V: 0.001 to 0.15%, W: 0.001 to 0.15%, Mo: 0 by mass% 0.001 to 0.45%, Cu: 0.001 to 0.45%, Ni: 0.001 to 0.45%, Ca: 0.0001 to 0.005%, Sb: 0.0001 to 0.10 The steel member according to [2], containing one or more selected from%.
[4] The steel member according to any one of [1] to [3], wherein the steel member is a welded steel pipe.
［５］前記［１］～［４］のいずれかに記載の鋼部材用の熱延鋼板であって、質量％で、Ｔｉを０．０３１～０．２００％含有し、組織中に０．００５％以上のＴｉが固溶Ｔｉとして存在する、鋼部材用の熱延鋼板。
［６］長手方向両端部である先端部および尾端部の板厚が、ともに長手方向中央部の板厚に比べて５～５０％厚い、［５］に記載の鋼部材用の熱延鋼板。 [5] A hot-rolled steel sheet for steel members as set forth in any one of [1] to [4] above, containing 0.031 to 0.200% Ti by mass% and having a structure of 0.001. A hot-rolled steel sheet for steel members in which 005% or more of Ti exists as solute Ti.
[6] The hot-rolled steel sheet for steel members according to [5], wherein the plate thicknesses at the tip and tail ends that are both ends in the longitudinal direction are both 5 to 50% thicker than the plate thickness at the center in the longitudinal direction. .
［７］前記［１］～［４］のいずれかに記載の鋼部材の製造方法であって、質量％で、Ｔｉを０．０３１～０．２００％含有し、組織中に０．００５％以上のＴｉが固溶Ｔｉとして存在する熱延鋼板に成形加工を施した後に、５５０℃を超え１０５０℃以下の温度に加熱した後、５５０～４００℃の温度域を１０℃／ｓ以上の平均冷却速度で冷却する熱処理を施す、鋼部材の製造方法。
［８］前記熱延鋼板を、質量％で、Ｔｉを０．０３１～０．２００％含有する鋼スラブを下記（１）式から計算される平衡固溶温度ＴＴｉよりも高い温度条件でスラブ抽出した後、ＴＴｉ－４００℃以上の温度で仕上げ圧延を終了し、ＴＴｉ－４００℃からＴＴｉ－５００℃までの温度域を１０℃／ｓ以上の平均冷却速度で冷却し、ＴＴｉ－５００℃以下の温度で巻き取って製造する、［７］に記載の鋼部材の製造方法。
ｌｏｇ（［Ｔｉ－Ｎ×４８÷１４］［Ｃ］）＝－７０００／（ＴＴｉ（℃）＋２７３）＋２．７５　・・・（１）
ただし、（１）式におけるＴｉ、Ｎ、Ｃは、鋼スラブ中のそれぞれの元素の含有量（質量％）である。 [7] The method for manufacturing a steel member according to any one of [1] to [4], wherein 0.031 to 0.200% of Ti is contained by mass and 0.005% in the structure. After forming the hot-rolled steel sheet in which Ti is present as solute Ti and heating it to a temperature exceeding 550 ° C. and not exceeding 1050 ° C., the temperature range of 550 to 400 ° C. is an average of 10 ° C./s or more. The manufacturing method of the steel member which performs the heat processing cooled at a cooling rate.
[8] A steel slab containing the hot rolled steel sheet in mass% and Ti in an amount of 0.031 to 0.200% under a temperature condition higher than the equilibrium solid solution temperature T Ti calculated from the following equation (1). after extraction, exit finish rolling at T Ti -400 ° C. or higher, the temperature range from T Ti -400 ° C. until T Ti -500 ° C. and cooled at 10 ° C. / s or more average cooling rate, T Ti The method for manufacturing a steel member according to [7], wherein the steel member is wound and manufactured at a temperature of −500 ° C. or lower.
log ([Ti−N × 48 ÷ 14] [C]) = − 7000 / (T Ti (° C.) + 273) +2.75 (1)
However, Ti, N, and C in the formula (1) are contents (mass%) of respective elements in the steel slab.
［９］前記［５］または［６］に記載の鋼部材用の熱延鋼板の製造方法であって、質量％で、Ｔｉを０．０３１～０．２００％含有する鋼スラブを、下記（１）式から計算される平衡固溶温度ＴＴｉよりも高い温度条件でスラブ抽出した後、ＴＴｉ－４００℃以上の温度で仕上げ圧延を終了し、ＴＴｉ－４００℃からＴＴｉ－５００℃までの温度域を１０℃／ｓ以上の平均冷却速度で冷却し、ＴＴｉ－５００℃以下の温度で巻き取る、鋼部材用の熱延鋼板の製造方法。
ｌｏｇ（［Ｔｉ－Ｎ×４８÷１４］［Ｃ］）＝－７０００／（ＴＴｉ（℃）＋２７３）＋２．７５　・・・（１）
ただし、（１）式におけるＴｉ、Ｎ、Ｃは、鋼スラブ中のそれぞれの元素の含有量（質量％）である。 [9] A method for producing a hot-rolled steel sheet for steel members according to [5] or [6], wherein a steel slab containing 0.031 to 0.200% Ti by mass% is 1) After slab extraction under a temperature condition higher than the equilibrium solid solution temperature T Ti calculated from the equation, finish rolling is finished at a temperature of T Ti −400 ° C. or more, and from T Ti −400 ° C. to T Ti −500 ° C. A method for producing a hot-rolled steel sheet for a steel member, wherein the temperature range is cooled at an average cooling rate of 10 ° C./s or more and wound at a temperature of T Ti −500 ° C. or less.
log ([Ti−N × 48 ÷ 14] [C]) = − 7000 / (T Ti (° C.) + 273) +2.75 (1)
However, Ti, N, and C in the formula (1) are contents (mass%) of respective elements in the steel slab.
　本発明によれば、塑性歪域の耐疲労特性に優れる鋼部材を提供できる。また、本発明の熱延鋼板は、前記鋼部材の素材として特に適する。
　本発明によれば、強度と塑性歪域での耐疲労特性という相反する特性を高度なレベルで両立できる鋼部材を提供できる。そのため、本発明の鋼部材としては、特に、高強度で塑性歪域の耐疲労特性を要求される、コイルドチュービング用溶接鋼管、ラインパイプ用溶接鋼管、自動車用構造部材用溶接鋼管が好適であり、なかでもコイルドチュービング用溶接鋼管が好適である。 ADVANTAGE OF THE INVENTION According to this invention, the steel member excellent in the fatigue resistance characteristic of a plastic strain area can be provided. The hot-rolled steel sheet of the present invention is particularly suitable as a material for the steel member.
ADVANTAGE OF THE INVENTION According to this invention, the steel member which can make compatible the characteristic which is the strength and the fatigue-resistant characteristic in a plastic strain area at a high level can be provided. Therefore, as the steel member of the present invention, a welded steel pipe for coiled tubing, a welded steel pipe for line pipes, and a welded steel pipe for automotive structural members, which are particularly required to have high strength and fatigue resistance in the plastic strain region, are suitable. In particular, a welded steel pipe for coiled tubing is suitable.
後熱処理により粒径２０ｎｍ以下の析出物として析出したＴｉ量と塑性歪域での疲労特性の関係を示す図である。It is a figure which shows the relationship between the amount of Ti precipitated as a precipitate with a particle size of 20 nm or less by post-heat treatment, and the fatigue characteristics in a plastic strain region.
（鋼部材）
　本発明の鋼部材は、特定温度加工条件で熱間圧延して製造した熱延鋼板（素材熱延鋼板）に成形加工を施した後、特定の条件で熱処理することで得られる。以下、素材熱延鋼板に成形加工を施した後に施す熱処理を「後熱処理」ともいう。
　まず、本発明の鋼部材の化学成分範囲の限定理由について説明する。なお、以下、組成における質量％は、単に％で示す。 (Steel member)
The steel member of the present invention is obtained by subjecting a hot-rolled steel sheet (raw material hot-rolled steel sheet) produced by hot rolling under specific temperature processing conditions to a heat treatment under specific conditions. Hereinafter, the heat treatment performed after forming the raw hot-rolled steel sheet is also referred to as “post-heat treatment”.
First, the reason for limiting the chemical component range of the steel member of the present invention will be described. Hereinafter, the mass% in the composition is simply represented by%.
　Ｔｉ：０．０３１～０．２００％
　Ｔｉは、熱間圧延工程で炭窒化物として析出し、熱間圧延工程での回復・再結晶の粒成長を抑制する。Ｔｉを含有することで、素材熱延鋼板の組織（ミクロ組織）中に所望の微細なフェライト相の粒径（１～５０μｍ）を得られる効果がある。この素材熱延鋼板段階でのミクロ組織の微細化は、その後の造管、部品成形等の成形加工（冷間加工）後に、熱処理を施した後のミクロ組織の微細化につながり、優れた塑性歪域での耐疲労特性が得られる。 Ti: 0.031 to 0.200%
Ti precipitates as carbonitride in the hot rolling process, and suppresses recovery / recrystallization grain growth in the hot rolling process. By containing Ti, there is an effect that a desired fine ferrite phase particle size (1 to 50 μm) can be obtained in the structure (microstructure) of the hot-rolled steel sheet. The refinement of the microstructure at the hot-rolled steel sheet stage leads to the refinement of the microstructure after heat treatment after subsequent forming (cold working) such as pipe forming and part forming, and excellent plasticity. Fatigue resistance in the strain range is obtained.
　Tanakaらは疲労サイクルによってすべり面上に転位が不可逆的にパイルアップし、このときに発生する応力が限界応力を超えると初期亀裂が発生するというモデルを提唱している（文献：K. Tanaka and T. Mura：J Appl Mech., Vol. 48, p.97-103 (1981)）。このモデルによれば、Ｇ：横弾性定数、Ｗｓ：単位面積あたりの破壊エネルギー、ν：ポアソン比、Δτ：すべり面上の分解剪断応力範囲、ｋ：すべり面上の転位の摩擦力など材料物性値、外力条件等が同じであれば、各結晶粒の疲労亀裂発生サイクルＮｃは、すべり面長さｄが短いほど、すなわち結晶粒径が小さいほど長くなる。
　このようなメカニズムにより、本発明の微細化されたミクロ組織材は、疲労き裂発生が遅れ、優れた塑性歪域での耐疲労特性を示すものと考えられる。 Tanaka et al. Proposed a model in which dislocations pile up irreversibly on the slip surface due to fatigue cycles, and when the stress generated at this time exceeds the critical stress, an initial crack occurs (reference: K. Tanaka and T. Mura: J Appl Mech., Vol. 48, p.97-103 (1981)). According to this model, G: transverse elastic constant, Ws: fracture energy per unit area, ν: Poisson's ratio, Δτ: decomposition shear stress range on the sliding surface, k: dislocation frictional force on the sliding surface, etc. If the value, the external force condition, and the like are the same, the fatigue crack generation cycle Nc of each crystal grain becomes longer as the slip surface length d is shorter, that is, as the crystal grain size is smaller.
Due to such a mechanism, it is considered that the refined microstructure material of the present invention is delayed in fatigue crack generation and exhibits excellent fatigue resistance characteristics in the plastic strain region.
　さらにＴｉは炭化物としてマトリクスを析出強化し、かつ固溶元素として固溶強化し、かつ焼入れ性向上元素として変態組織強化を強めることで、造管、部品成形等の成形加工後に熱処理を施した後の強度が向上し、疲労強度を著しく向上させる必須元素である。こうした効果は、Ｔｉ含有量が０．０３１～０．２００％の範囲にあるときに得られ、Ｔｉ含有量が前記範囲の下限値未満であると、後述する素材熱延鋼板の段階で０．００５％以上のＴｉが固溶Ｔｉとして存在し、成形加工した後の熱処理によって、０．００５％以上のＴｉを粒径２０ｎｍ以下の微細な析出物として析出させることができず、上記効果を得ることができない。一方、Ｔｉ含有量が前記範囲の上限値を超えると、粗大なＴｉＮの生成によって耐疲労特性が低下する。そのため、Ｔｉ含有量は０．０３１～０．２００％の範囲とする。Ｔｉ含有量は、好ましくは０．１２０％超である。また、Ｔｉ含有量は、好ましくは０．１５０％以下である。 In addition, Ti is strengthened by precipitation strengthening the matrix as a carbide, solid solution strengthening as a solid solution element, and strengthening transformation structure strengthening as a hardenability improving element, and after heat treatment after forming processing such as pipe making and part forming This is an essential element that improves the strength of the steel and significantly improves the fatigue strength. Such an effect is obtained when the Ti content is in the range of 0.031 to 0.200%, and when the Ti content is less than the lower limit of the above range, the effect is 0. 005% or more of Ti exists as solute Ti, and the heat treatment after the molding process cannot precipitate 0.005% or more of Ti as fine precipitates having a particle diameter of 20 nm or less, and the above effect is obtained. I can't. On the other hand, if the Ti content exceeds the upper limit of the above range, the fatigue resistance is reduced due to the formation of coarse TiN. Therefore, the Ti content is in the range of 0.031 to 0.200%. The Ti content is preferably more than 0.120%. Further, the Ti content is preferably 0.150% or less.
　本発明の鋼部材の組織中には、０．００５％以上のＴｉが粒径２０ｎｍ以下の析出物として析出している。
　本発明者らは、本発明のように熱延鋼板を素材とし、さらに造管あるいは部品成形等の成形加工後に施される熱処理（後熱処理）の後に、塑性歪域での耐疲労特性が必要とされる場合、後熱処理によって、０．００５％以上のＴｉを粒径２０ｎｍ以下の微細な析出物として析出させることで、格段に優れた塑性歪域での耐疲労特性が得られることを知見した。図１に、後熱処理により粒径２０ｎｍ以下の微細な析出物として析出したＴｉ量（質量％）と塑性歪域での疲労特性の関係を示す。後熱処理により粒径２０ｎｍ以下の微細な析出物として析出したＴｉ量が０．００５％以上になると、引張モード、ひずみ制御モード、ひずみ比＝０、全ひずみ範囲２．０％の条件で引張疲労試験を行った場合の破断までの繰り返し数が１０００回以上となり、優れた塑性歪域での耐疲労特性が得られる。 In the structure of the steel member of the present invention, 0.005% or more of Ti is precipitated as a precipitate having a particle size of 20 nm or less.
The present inventors need a fatigue resistance property in a plastic strain region after a heat treatment (post heat treatment) performed after a forming process such as pipe making or part forming using a hot-rolled steel sheet as in the present invention. In that case, it is found that by applying post-heat treatment, 0.005% or more of Ti is precipitated as fine precipitates having a particle size of 20 nm or less, so that excellent fatigue resistance characteristics in the plastic strain region can be obtained. did. FIG. 1 shows the relationship between the amount of Ti (mass%) precipitated as a fine precipitate having a particle size of 20 nm or less by post-heat treatment and the fatigue characteristics in the plastic strain region. When the amount of Ti deposited as a fine precipitate having a particle size of 20 nm or less by post-heat treatment reaches 0.005% or more, tensile fatigue occurs under the conditions of tensile mode, strain control mode, strain ratio = 0, and total strain range of 2.0%. When the test is performed, the number of repetitions until breakage is 1000 times or more, and excellent fatigue resistance characteristics in the plastic strain region can be obtained.
　次に、本発明の鋼部材が有する好適な組成について説明する。 Next, the preferred composition of the steel member of the present invention will be described.
　Ｃ：０．１９～０．５０％
　本発明において、Ｃは、特定の条件で後熱処理することで、高い強度を確保させ、さらに後熱処理時にＴｉと結合し、特に表層部おいて微細析出物を析出させ塑性歪域での耐疲労特性を向上させる元素である。Ｃの含有量が０．１９％未満では、この所望の強度（ＹＳ≧７７０ＭＰａ）と塑性歪域での耐疲労特性を得られにくくなる。一方、Ｃの含有量が０．５０％を超えると、鋼部材、例えば鋼管の靱性、溶接性が確保できなくなるため、これを上限とする。なお、さらに好ましくは、Ｃの含有量は０．２８％超である。また、さらに好ましくは、Ｃの含有量は０．３０％以下である。 C: 0.19 to 0.50%
In the present invention, C is post-heat-treated under specific conditions to ensure high strength, and further binds to Ti during post-heat treatment, in particular, precipitates fine precipitates in the surface layer portion to cause fatigue resistance in the plastic strain region. It is an element that improves the characteristics. When the C content is less than 0.19%, it becomes difficult to obtain the desired strength (YS ≧ 770 MPa) and fatigue resistance in the plastic strain region. On the other hand, if the C content exceeds 0.50%, the toughness and weldability of a steel member, for example, a steel pipe, cannot be ensured, so this is the upper limit. More preferably, the C content is more than 0.28%. More preferably, the C content is 0.30% or less.
　Ｓｉ：０．００２～１．５％
　Ｓｉは、固溶強化により所望の強度を確保しつつ、塑性歪域での耐疲労特性を向上させる元素である。Ｓｉの含有量が０．００２％未満では強度が不足する。一方、１．５％を超える含有は、溶接性が劣化する。従ってＳｉの含有量は０．００２～１．５％に限定することが好ましい。なお、さらに好ましくは、Ｓｉの含有量は０．０５％以上である。また、さらに好ましくは、Ｓｉの含有量は０．３５％以下である。 Si: 0.002 to 1.5%
Si is an element that improves fatigue resistance in the plastic strain region while ensuring a desired strength by solid solution strengthening. If the Si content is less than 0.002%, the strength is insufficient. On the other hand, if the content exceeds 1.5%, weldability deteriorates. Therefore, the Si content is preferably limited to 0.002 to 1.5%. More preferably, the Si content is 0.05% or more. More preferably, the Si content is 0.35% or less.
　Ｍｎ：０．４～２．５％
　Ｍｎは、後熱処理時に低温変態強化により所望の強度を確保させ、塑性歪域での耐疲労特性を向上させる働きがある。Ｍｎの含有量が０．４％未満では、この効果が十分に発現せず、一方、Ｍｎの含有量が２．５％を超えると溶接性が劣化する。従ってＭｎの含有量は０．４～２．５％に限定することが好ましい。なお、さらに好ましくは、Ｍｎの含有量は１．０９％以上である。また、さらに好ましくは、Ｍｎの含有量は１．９９％以下である。 Mn: 0.4 to 2.5%
Mn has a function of ensuring a desired strength by strengthening at low temperature during post-heat treatment and improving fatigue resistance in the plastic strain region. If the Mn content is less than 0.4%, this effect is not sufficiently exhibited. On the other hand, if the Mn content exceeds 2.5%, the weldability deteriorates. Therefore, the Mn content is preferably limited to 0.4 to 2.5%. More preferably, the Mn content is 1.09% or more. More preferably, the Mn content is 1.99% or less.
　Ａｌ：０．０１～０．１９％
　Ａｌは、製鋼時の脱酸元素であるとともに、熱間圧延工程でのオーステナイト粒の成長を抑制し、結晶粒を微細とし、後熱処理後に所望のフェライト粒径（１～５０μｍ）を得られ、塑性歪域での耐疲労特性を向上させる働きがある。Ａｌの含有量が０．０１％未満ではこれらの効果が得られずフェライト粒径が粗大化し、一方、Ａｌの含有量が０．１９％を超えると溶接性が低下するともに、酸化物系介在物の増大により耐疲労特性が低下する傾向となる。なお、さらに好ましくは、Ａｌの含有量は０．０４１％以上である。また、さらに好ましくは、Ａｌの含有量は０．０８０％以下である。 Al: 0.01 to 0.19%
Al is a deoxidizing element during steel making, suppresses the growth of austenite grains in the hot rolling process, makes the crystal grains fine, and obtains a desired ferrite grain size (1 to 50 μm) after post-heat treatment, It has the function of improving fatigue resistance in the plastic strain region. If the Al content is less than 0.01%, these effects cannot be obtained, and the ferrite grain size becomes coarse. On the other hand, if the Al content exceeds 0.19%, the weldability deteriorates and the oxide type intervening The fatigue resistance tends to decrease due to the increase in the number of objects. More preferably, the Al content is 0.041% or more. More preferably, the Al content is 0.080% or less.
　Ｃｒ：０．００１～０．９０％
　Ｃｒは、後熱処理時に低温変態強化により所望の強度を確保させ、塑性歪域での耐疲労特性を向上させる働きがある。Ｃｒの含有量が０．００１％未満では、この効果が十分に発現せず、一方、Ｃｒの含有量が０．９０％を超えると溶接性が劣化する。従って、Ｃｒの含有量は０．００１～０．９０％に限定することが好ましい。なお、さらに好ましくは、Ｃｒの含有量は０．００１～０．１９％である。 Cr: 0.001 to 0.90%
Cr has a function of securing a desired strength by strengthening at low temperature transformation during post-heat treatment and improving fatigue resistance in a plastic strain region. If the Cr content is less than 0.001%, this effect is not sufficiently exhibited. On the other hand, if the Cr content exceeds 0.90%, the weldability deteriorates. Therefore, the Cr content is preferably limited to 0.001 to 0.90%. More preferably, the Cr content is 0.001 to 0.19%.
　Ｂ：０．０００１～０．００５０％
　Ｂは、後熱処理時に低温変態強化により所望の強度を確保させ、塑性歪域での耐疲労特性を向上させる働きがある。Ｂの含有量が０．０００１％未満では、この効果が十分に発現せず、一方、Ｂの含有量が０．００５０％を超えると耐疲労特性が低下する傾向となる。従って、Ｂの含有量は０．０００１～０．００５０％に限定することが好ましい。なお、さらに好ましくは、Ｂの含有量は０．０００５％以上である。また、さらに好ましくは、Ｂの含有量は０．００３５％以下である。 B: 0.0001 to 0.0050%
B has a function of ensuring a desired strength by strengthening at low temperature transformation during post-heat treatment and improving fatigue resistance in a plastic strain region. If the B content is less than 0.0001%, this effect is not sufficiently exhibited. On the other hand, if the B content exceeds 0.0050%, the fatigue resistance tends to decrease. Therefore, the B content is preferably limited to 0.0001 to 0.0050%. More preferably, the B content is 0.0005% or more. More preferably, the B content is 0.0035% or less.
　Ｐ：０．０１９％以下（０％を含む）
　Ｐは、Ｍｎとの凝固共偏析を介し、塑性歪域での耐疲労特性を低下させるとともに電縫溶接性を劣化させる。Ｐの含有量が０．０１９％を超えると悪影響が顕著となるため、０．０１９％を上限とすることが好ましい。 P: 0.019% or less (including 0%)
P deteriorates fatigue resistance in the plastic strain region and deteriorates electroweldability through solidification co-segregation with Mn. If the P content exceeds 0.019%, the adverse effect becomes remarkable, so 0.019% is preferable as the upper limit.
　Ｓ：０．０１５％以下（０％を含む）
　Ｓは、ＭｎＳなどとして鋼中介在物として存在し、塑性歪域での疲労亀裂の起点として耐疲労特性を低下させる。Ｓの含有量が０．０１５％を超えるとこの悪影響が顕著となるため、０．０１５％を上限とすることが好ましい。なお、さらに好ましくはＳの含有量は０．００５％以下である。 S: 0.015% or less (including 0%)
S exists as an inclusion in steel as MnS or the like, and lowers fatigue resistance as a starting point of fatigue cracks in the plastic strain region. When the S content exceeds 0.015%, this adverse effect becomes significant, so it is preferable to set the upper limit at 0.015%. More preferably, the S content is 0.005% or less.
　Ｎ：０．００８％以下（０％を含む）
　Ｎは、ＴｉとＴｉＮを形成し、粗大な析出物として析出し、固溶Ｔｉを消費する。こうしてＮは、Ｔｉ添加によって素材熱延鋼板の段階で０．００５％以上のＴｉを固溶Ｔｉとして存在させ、成形加工後の熱処理によって、０．００５％以上のＴｉを粒径２０ｎｍ以下の微細な析出物として析出し、格段に優れた塑性歪域での耐疲労特性が得られる効果を低下させる。Ｎの含有量が０．００８％を超えるとこの悪影響が顕著となるため、０．００８％を上限とすることが好ましい。なお、さらに好ましくは、Ｎの含有量は０．００４９％以下である。 N: 0.008% or less (including 0%)
N forms Ti and TiN, precipitates as coarse precipitates, and consumes solid solution Ti. Thus, N is added in the form of hot-rolled steel sheet at the raw material stage so that 0.005% or more of Ti is present as solute Ti, and heat treatment after forming is performed so that 0.005% or more of Ti is finely grained with a particle size of 20 nm or less. It precipitates as a good precipitate and reduces the effect of obtaining excellent fatigue resistance characteristics in the plastic strain region. If the N content exceeds 0.008%, this adverse effect becomes significant, so it is preferable to set the upper limit to 0.008%. More preferably, the N content is 0.0049% or less.
　Ｏ：０．００３％以下（０％を含む）
　Ｏは、酸化物系介在物として存在し、鋼の耐疲労特性を低下させる。Ｏの含有量が０．００３％を超えるとこの悪影響が顕著となるため、０．００３％を上限とすることが好ましい。なお、さらに好ましくは、Ｏの含有量は０．００２％以下である。 O: 0.003% or less (including 0%)
O exists as oxide inclusions and reduces the fatigue resistance of steel. If the content of O exceeds 0.003%, this adverse effect becomes remarkable, so 0.003% is preferably set as the upper limit. More preferably, the O content is 0.002% or less.
　Ｓｎ：０．１０％以下（０％を含む）
　Ｓｎは、固溶元素として存在し、鋼の熱間延性を低下させる。Ｓｎの含有量が０．１０％を超えるとこの悪影響が顕著となるため、０．１０％を上限とすることが好ましい。なお、さらに好ましくは、Ｓｎの含有量は０．０３％以下である。 Sn: 0.10% or less (including 0%)
Sn exists as a solid solution element and reduces the hot ductility of steel. If the Sn content exceeds 0.10%, this adverse effect becomes significant, so it is preferable to set the upper limit to 0.10%. More preferably, the Sn content is 0.03% or less.
　残部はＦｅおよび不可避的不純物である。本発明では、さらに、本発明の効果を向上させること等を目的に、つぎの元素を添加することができる。 The balance is Fe and inevitable impurities. In the present invention, the following elements can be further added for the purpose of improving the effects of the present invention.
　Ｎｂ：０．００１～０．１５％
　Ｎｂは、炭化物として析出し、熱間圧延工程での回復・再結晶の粒成長を抑制し、所望のフェライト粒径（１～５０μｍ）を得られる効果があり必要に応じて含有できる。Ｎｂの含有量が０．００１％未満ではこれらの効果が得られない。一方、Ｎｂの含有量が０．１５％を超えると、熱間圧延時の歪誘起析出によって表層部に粗大な析出物が析出し、表層部の微細析出物が減少し、塑性歪域での耐疲労特性が低下するため、０．１５％を上限とする。そのためＮｂを含有する場合には、Ｎｂの含有量を０．００１～０．１５％とする。なお、さらに好ましくは、Ｎｂの含有量は０．００１～０．００９％である。 Nb: 0.001 to 0.15%
Nb precipitates as a carbide, suppresses recovery / recrystallization grain growth in the hot rolling process, and has the effect of obtaining a desired ferrite grain size (1 to 50 μm), and can be contained as needed. If the Nb content is less than 0.001%, these effects cannot be obtained. On the other hand, when the content of Nb exceeds 0.15%, coarse precipitates are deposited on the surface layer portion due to strain-induced precipitation during hot rolling, and fine precipitates on the surface layer portion are reduced. Since fatigue resistance is reduced, the upper limit is made 0.15%. Therefore, when Nb is contained, the Nb content is set to 0.001 to 0.15%. More preferably, the Nb content is 0.001 to 0.009%.
　Ｖ：０．００１～０．１５％
　Ｖは、炭化物として析出し、熱間圧延工程での回復・再結晶の粒成長を抑制し、所望のフェライト粒径（１～５０μｍ）を得られる効果があり必要に応じて含有できる。Ｖの含有量が０．００１％未満ではこれらの効果が得られない。一方、Ｖの含有量が０．１５％を超えると、熱間圧延時の歪誘起析出によって表層部に粗大な析出物が析出し、表層部の微細析出物が減少し、塑性歪域での耐疲労特性が低下するため０．１５％を上限とする。そのためＶを含有する場合には、Ｖの含有量を０．００１～０．１５％とする。なお、さらに好ましくはＶの含有量は０．００１～０．０４９％である。 V: 0.001 to 0.15%
V precipitates as a carbide, suppresses recovery / recrystallization grain growth in the hot rolling process, and has the effect of obtaining a desired ferrite grain size (1 to 50 μm), and can be contained as necessary. If the V content is less than 0.001%, these effects cannot be obtained. On the other hand, when the content of V exceeds 0.15%, coarse precipitates are deposited on the surface layer portion due to strain-induced precipitation during hot rolling, and fine precipitates on the surface layer portion are reduced, and in the plastic strain region. The upper limit is 0.15% because the fatigue resistance is reduced. Therefore, when V is contained, the content of V is set to 0.001 to 0.15%. More preferably, the V content is 0.001 to 0.049%.
　Ｗ：０．００１～０．１５％
　Ｗは、炭化物として析出し、熱間圧延工程での回復・再結晶の粒成長を抑制し、所望のフェライト粒径（１～５０μｍ）を得られる効果を補完する働きがあり必要に応じて含有できる。Ｗの含有量が０．００１％未満ではこれらの効果が得られない。一方、Ｗの含有量が０．１５％を超えると、熱間圧延時の歪誘起析出によって表層部に粗大な析出物が析出し、表層部の微細析出物が減少し、塑性歪域での耐疲労特性が低下するため０．１５％を上限とする。そのためＷを含有する場合には、Ｗの含有量を０．００１～０．１５％とする。なお、さらに好ましくは、Ｗの含有量は０．００１～０．０４９％である。 W: 0.001 to 0.15%
W precipitates as carbide, suppresses recovery / recrystallization grain growth in the hot rolling process, supplements the effect of obtaining the desired ferrite grain size (1 to 50 μm), and is contained if necessary it can. If the W content is less than 0.001%, these effects cannot be obtained. On the other hand, when the content of W exceeds 0.15%, coarse precipitates are deposited on the surface layer portion due to strain-induced precipitation during hot rolling, and fine precipitates on the surface layer portion are reduced. The upper limit is 0.15% because the fatigue resistance is reduced. Therefore, when W is contained, the W content is set to 0.001 to 0.15%. More preferably, the W content is 0.001 to 0.049%.
　Ｍｏ：０．００１～０．４５％
　Ｍｏは、後熱処理時に低温変態強化或いは析出強化により所望の強度を確保させ、塑性歪域での耐疲労特性を向上させる効果を補完する働きがあり必要に応じて含有できる。Ｍｏの含有量が０．００１％未満では、この効果が発現せず、一方、Ｍｏの含有量が０．４５％を超えると溶接性が劣化する。従って、Ｍｏを含有する場合には、Ｍｏの含有量を０．００１～０．４５％とする。なお、さらに好ましくは、Ｍｏの含有量は０．００１～０．３０％である。 Mo: 0.001 to 0.45%
Mo has a function of securing a desired strength by low-temperature transformation strengthening or precipitation strengthening during post-heat treatment and complementing the effect of improving fatigue resistance in the plastic strain region, and can be contained as necessary. If the Mo content is less than 0.001%, this effect does not appear. On the other hand, if the Mo content exceeds 0.45%, the weldability deteriorates. Therefore, when Mo is contained, the Mo content is set to 0.001 to 0.45%. More preferably, the Mo content is 0.001 to 0.30%.
　Ｃｕ：０．００１～０．４５％、Ｎｉ：０．００１～０．４５％
　Ｃｕ、Ｎｉは、Ｍｎの疲労強度を向上させる効果を補完する働きがある元素であると同時に、鋼材の耐食性を高める効果があり、必要に応じてＣｕ、Ｎｉをそれぞれ含有できる。これら効果はＣｕ、Ｎｉそれぞれ０．００１％以上の含有で発現するが、Ｃｕ、Ｎｉそれぞれ０．４５％を超える含有は溶接性を低下させるために、それぞれ０．４５％を上限とする。そのためＣｕを含有する場合には、Ｃｕの含有量を０．００１～０．４５％とする。また、Ｎｉを含有する場合には、Ｎｉの含有量を０．００１～０．４５％とする。なお、さらに好ましくはいずれの元素も０．３５％以下である。 Cu: 0.001 to 0.45%, Ni: 0.001 to 0.45%
Cu and Ni are elements that have a function of complementing the effect of improving the fatigue strength of Mn, and at the same time, have the effect of increasing the corrosion resistance of the steel material, and can contain Cu and Ni as needed. These effects are manifested when the content of Cu and Ni is 0.001% or more. However, when the content exceeds 0.45% for Cu and Ni, the upper limit is 0.45% for lowering weldability. Therefore, when Cu is contained, the Cu content is set to 0.001 to 0.45%. When Ni is contained, the Ni content is set to 0.001 to 0.45%. More preferably, any element is 0.35% or less.
　Ｃａ：０．０００１～０．００５％
　Ｃａは、展伸したＭｎＳを粒状のＣａ（Ａｌ）Ｓ（Ｏ）とする所謂形態制御効果があり、疲労亀裂発生を抑制し、耐疲労特性を向上させる効果があり、必要に応じて含有できる。この効果は０．０００１％以上の含有で発現するが、０．００５％を超える含有は、非金属介在物の増大によってかえって耐疲労特性が低下するために０．００５％を上限とする。そのためＣａを含有する場合には、Ｃａの含有量を０．０００１～０．００５％とする。 Ca: 0.0001 to 0.005%
Ca has a so-called form control effect in which expanded MnS is granular Ca (Al) S (O), has the effect of suppressing fatigue cracking and improving fatigue resistance, and can be contained if necessary. . This effect is manifested with a content of 0.0001% or more. However, the content exceeding 0.005% is limited to 0.005% because the fatigue resistance is lowered by the increase of nonmetallic inclusions. Therefore, when Ca is contained, the content of Ca is set to 0.0001 to 0.005%.
　Ｓｂ：０．０００１～０．１０％
　Ｓｂは、表面に優先的に偏析し、熱間圧延工程、或いは後熱処理工程での雰囲気からのＮの侵入を抑制し、ＢＮの形成によるＢの添加効果の減少を抑制する働きがあり、必要に応じて含有できる。この効果は、０．０００１％以上の含有で発現するが、０．１０％を超えて含有しても効果が飽和するために０．１０％を上限とする。そのためＳｂを含有する場合には、Ｓｂの含有量を０．０００１～０．１０％とする。なお、さらに好ましくは、Ｓｂの含有量は０．０００１～０．０３０％である。 Sb: 0.0001 to 0.10%
Sb preferentially segregates on the surface, suppresses the intrusion of N from the atmosphere in the hot rolling process or the post heat treatment process, and functions to suppress a decrease in the effect of adding B due to the formation of BN. Depending on the content. This effect is manifested at a content of 0.0001% or more, but even if it exceeds 0.10%, the effect is saturated, so 0.10% is made the upper limit. Therefore, when Sb is contained, the Sb content is set to 0.0001 to 0.10%. More preferably, the Sb content is 0.0001 to 0.030%.
　また、本発明の鋼部材は、後熱処理後の表面から板厚方向２００μｍまでのフェライト相の平均結晶粒径が１～５０μｍであり、表面から板厚方向２００μｍまでのフェライト相中に粒径１．０～２０ｎｍのＴｉ炭化物が析出してなる組織を有し、表面から板厚方向２００μｍまでの平均硬度と、中心偏析部を除く板厚中心近傍の平均硬度の差（絶対値）が、ビッカース硬さ（ＨＶ）で、ΔＨＶ５０ポイント以下であることが望ましい。 In the steel member of the present invention, the average crystal grain size of the ferrite phase from the surface after post heat treatment to the plate thickness direction of 200 μm is 1 to 50 μm, and the particle size of 1 in the ferrite phase from the surface to the plate thickness direction of 200 μm. The difference between the average hardness from the surface to the thickness direction of 200 μm and the average hardness in the vicinity of the thickness center excluding the central segregation part (absolute value) has a structure in which Ti carbide of 0.0 to 20 nm is precipitated. The hardness (HV) is desirably ΔHV 50 points or less.
　鋼部材のミクロ組織、析出物の析出状態、並びに断面硬度は、優れた塑性歪域での耐疲労特性を確保する上で重要である。後熱処理後の表面から板厚方向２００μｍまでのフェライト相の平均結晶粒径が５０μｍ超えでは、初期疲労き裂が早く、大きく発生し所望の塑性歪域での耐疲労特性が確保しにくくなる。一方、後熱処理後にフェライト相の平均結晶粒径を１μｍ未満とすることは工業的、経済的に難しいためこれを下限とした。 The microstructure of the steel member, the precipitation state of the precipitates, and the cross-sectional hardness are important for ensuring fatigue resistance in an excellent plastic strain region. If the average crystal grain size of the ferrite phase from the surface after post-heat treatment to the plate thickness direction of 200 μm exceeds 50 μm, the initial fatigue cracks are early and large, making it difficult to ensure fatigue resistance characteristics in a desired plastic strain region. On the other hand, since it is difficult industrially and economically to make the average crystal grain size of the ferrite phase less than 1 μm after the post heat treatment, this is set as the lower limit.
　なお、ここでいうフェライト相とは体心立方格子の母相鉄を謂い、ポリゴナルフェライト、アシキュラーフェライト、ウィッドマンステッテンフェライト、ベイニティックフェライト、ベイナイト、低炭素（Ｃ含有量１％以下）マルテンサイト組織を含むものとする。なお、フェライト相以外の第二相としては、オーステナイト、カーバイド、パーライト、高炭素マルテンサイト（Ｃ含有量１％超え）が挙げられる。
　本発明の鋼部材の組織は、上記フェライト相を主相とすることが好ましい。ここで、主相とは、体積率で、５１％以上占有する相をいい、８０％以上が好ましく、１００％であってもよい。 The term “ferrite phase” as used herein refers to the body phase iron of a body-centered cubic lattice, so-called polygonal ferrite, acicular ferrite, Widmanstatten ferrite, bainitic ferrite, bainite, and low carbon (C content of 1% or less). Includes a martensite organization. Examples of the second phase other than the ferrite phase include austenite, carbide, pearlite, and high carbon martensite (C content exceeding 1%).
The structure of the steel member of the present invention preferably has the ferrite phase as a main phase. Here, the main phase refers to a phase occupying 51% or more by volume ratio, preferably 80% or more, and may be 100%.
　また、表面から板厚方向２００μｍまでのフェライト相中のＴｉ炭化物寸法は、表面硬度を確保し、高い塑性歪域での耐疲労特性を確保するために重要である。表面から板厚方向２００μｍまでのフェライト相中に１．０～２０ｎｍのＴｉ炭化物が析出することで、疲労初期き裂の発生が抑制され、またその寸法が小さくなり優れた塑性歪域での耐疲労特性をより高めることができる。なお、１．０～２０ｎｍのＴｉ炭化物の析出量はここでは特に定めない。また、１．０～２０ｎｍのＴｉ炭化物以外に、寸法の異なるＴｉ炭化物が析出していることも許容する。 Also, the Ti carbide dimension in the ferrite phase from the surface to the plate thickness direction of 200 μm is important for ensuring the surface hardness and the fatigue resistance in a high plastic strain region. The precipitation of 1.0 to 20 nm Ti carbide in the ferrite phase from the surface to the thickness direction of 200 μm suppresses the occurrence of initial fatigue cracks, reduces the size, and provides excellent resistance to plastic strain. The fatigue characteristics can be further improved. Here, the precipitation amount of Ti carbide of 1.0 to 20 nm is not particularly defined here. In addition to Ti carbide having a thickness of 1.0 to 20 nm, it is allowed to deposit Ti carbides having different dimensions.
　表面から板厚方向２００μｍまでの平均硬度と、中心偏析部を除く板厚中心近傍の平均硬度の差がΔＨＶ５０ポイント以下であることは、優れた塑性歪域での耐疲労特性を確保する上で重要である。表面から板厚方向２００μｍまでの平均硬度と、中心偏析部を除く板厚中心近傍の平均硬度の差がΔＨＶ５０ポイントを超えると、初期疲労き裂が早く、大きく発生し、所望の塑性歪域での耐疲労特性を確保しにくくなる。このため、表面から板厚方向２００μｍまでの平均硬度と、中心偏析部を除く板厚中心近傍の平均硬度の差がΔＨＶ５０ポイント以下であることが望ましい。 The difference between the average hardness from the surface to the thickness direction of 200 μm and the average hardness in the vicinity of the thickness center excluding the center segregation part is ΔHV50 points or less, in order to ensure excellent fatigue resistance in the plastic strain region. is important. When the difference between the average hardness from the surface to the thickness direction of 200 μm and the average hardness in the vicinity of the center of the plate thickness excluding the center segregation part exceeds ΔHV50 points, the initial fatigue cracks occur quickly and greatly, and in the desired plastic strain region It is difficult to ensure the fatigue resistance characteristics. For this reason, it is desirable that the difference between the average hardness from the surface to the plate thickness direction of 200 μm and the average hardness in the vicinity of the plate thickness center excluding the center segregation portion is ΔHV 50 points or less.
　なお、表面から板厚方向２００μｍまでの平均硬度と、中心偏析部を除く板厚中心近傍の平均硬度の差は、板厚方向５０～２００μｍの間を板厚方向に２５μｍピッチでマイクロビッカース硬度を荷重０．１ｋｇｆで測定し(HV(0.1))、７点を平均した値HV(0.1)Sと、板厚中心部を中心に中心偏析部を避け、板厚方向に２５μｍピッチでHV(0.1)を７点測定し平均した値HV(0.1)Cの差、HV(0.1)C－HV(0.1)Sとして測定した。 The difference between the average hardness from the surface to the thickness direction of 200 μm and the average hardness in the vicinity of the thickness center excluding the center segregation part is the micro Vickers hardness at a pitch of 25 μm in the thickness direction between 50 and 200 μm in the thickness direction. Measured at a load of 0.1 kgf (HV (0.1)), averaged 7 points HV (0.1) S , avoiding the center segregation around the center of the plate thickness, HV (0.1 average value HV (0.1) the difference in C) was measured 7 points was measured as HV (0.1) C -HV (0.1 ) S.
（素材熱延鋼板）
　本発明の鋼部材用の熱延鋼板（素材熱延鋼板）は、本発明の鋼部材を得るために特に好適なものである。
　本発明の素材熱延鋼板は、質量％で、Ｔｉを０．０３１～０．２００％含有し、組織中に０．００５％以上のＴｉが固溶Ｔｉとして存在する。これにより、成形加工後、所定の熱処理を施した後に、鋼部材の組織中に０．００５％以上のＴｉを粒径２０ｎｍ以下の微細な析出物として析出させることができ、塑性歪域での耐疲労特性に優れ、さらに強度特性にも優れる鋼部材を得ることができる。 (Material hot-rolled steel sheet)
The hot-rolled steel sheet (raw material hot-rolled steel sheet) for steel members of the present invention is particularly suitable for obtaining the steel member of the present invention.
The material hot-rolled steel sheet of the present invention contains 0.031 to 0.200% Ti by mass%, and 0.005% or more of Ti exists in the structure as solute Ti. Thereby, after performing a predetermined heat treatment after forming, 0.005% or more of Ti can be precipitated as fine precipitates having a particle size of 20 nm or less in the structure of the steel member, and in the plastic strain region It is possible to obtain a steel member that is excellent in fatigue resistance and also excellent in strength characteristics.
　本発明の素材熱延鋼板の有する組成は、上記鋼部材の有する組成と同様である。
　また、本発明の素材熱延鋼板は、長手方向両端部である先端部および尾端部の板厚が、ともに長手方向両端部以外の中間部（長手方向中央部）の板厚に比べて５～５０％厚いことが好ましい。このことにより、コイルドチュービングのように、素材熱延鋼板を所定の幅にスリットした後、長手方向に溶接で繋いで用いる場合の、溶接部の塑性歪域での耐疲労特性が向上する効果が高められる。 The composition of the material hot-rolled steel sheet of the present invention is the same as the composition of the steel member.
In the hot-rolled steel sheet of the present invention, the thickness of the tip and tail ends, which are both ends in the longitudinal direction, is 5 compared to the thickness of the intermediate portion (longitudinal central portion) other than both ends in the longitudinal direction. Preferably it is ~ 50% thick. As a result, the effect of improving the fatigue resistance in the plastic strain region of the welded part when slitting the raw hot-rolled steel sheet to a predetermined width and connecting it by welding in the longitudinal direction as in coiled tubing Is increased.
（製造方法）
　次に、本発明の鋼部材と、その素材となる熱延鋼板の製造方法について説明する。なお、以下の説明において、特に断らない限り、温度は鋼スラブ等の表面温度とする。
　本発明では、上記した組成を有する鋼を鋳造した鋼スラブを出発素材とする。出発素材の製造方法は、とくに限定されず、例えば、上記した組成の溶鋼を転炉等の常用の溶製方法で溶製し、連続鋳造法等の通常の鋳造方法で鋼スラブとする方法等が挙げられる。 (Production method)
Next, the steel member of this invention and the manufacturing method of the hot rolled steel plate used as the raw material are demonstrated. In the following description, unless otherwise specified, the temperature is the surface temperature of a steel slab or the like.
In the present invention, a steel slab obtained by casting steel having the above composition is used as a starting material. The production method of the starting material is not particularly limited. For example, the molten steel having the above composition is melted by a conventional melting method such as a converter, and a steel slab is obtained by a normal casting method such as a continuous casting method. Is mentioned.
　まず、本発明の鋼部材の素材となる熱延鋼板（素材熱延鋼板）の製造方法について説明する。
　本発明の素材熱延鋼板は、Ｔｉを０．０３１～０．２００％含有する鋼スラブを所定の条件で熱間圧延することで製造できる。 First, the manufacturing method of the hot-rolled steel plate (raw material hot-rolled steel plate) used as the raw material of the steel member of this invention is demonstrated.
The material hot-rolled steel sheet of the present invention can be manufactured by hot rolling a steel slab containing 0.031 to 0.200% Ti under predetermined conditions.
　ｌｏｇ（［Ｔｉ－Ｎ×４８÷１４］［Ｃ］）＝－７０００／（ＴＴｉ（℃）＋２７３）＋２．７５から計算される平衡固溶温度ＴＴｉよりも高い温度条件でスラブ抽出
　熱間圧延工程におけるスラブ抽出温度は鋼中のＴｉの再固溶、析出状況を通じて、熱間圧延後の析出物サイズ、固溶Ｔｉ量に影響を及ぼし、後熱処理後に良好な耐疲労特性を確保するために重要である。抽出温度が下記（１）式から計算される平衡固溶温度ＴＴｉ以下であると、連続鋳造時に析出した粗大なＴｉが未固溶炭窒化物として残存し、素材熱延鋼板の段階で固溶Ｔｉ量が０．００５％未満となり、後熱処理後に格段に優れた塑性歪域での耐疲労特性が得られない。下記（１）式から計算される平衡固溶温度ＴＴｉよりも高い温度条件でスラブ抽出することで、素材熱延鋼板の段階で０．００５％以上のＴｉが固溶Ｔｉとして存在し、成形加工後の熱処理によって、０．００５％以上のＴｉを粒径２０ｎｍ以下の微細な析出物として析出させることができ、格段に優れた塑性歪域での耐疲労特性が得られる。なお、さらに好ましくは結晶粒径の粗大化防止の観点から、スラブ抽出温度は１６２０Ｋ以下であることが好ましく、Ｔｉの固溶状態の均一性と十分な固溶時間の確保の観点からスラブの均熱時間（平衡固溶温度ＴＴｉよりも高い温度でスラブを保持する時間）は１０ｍｉｎ以上であることが好ましい。 log ([Ti-N × 48 ÷ 14] [C]) = − 7000 / (T Ti (° C.) + 273) +2.75 Equilibrium solution temperature T slab extracted at a temperature higher than T Ti The slab extraction temperature in the rolling process affects the precipitate size after hot rolling and the amount of solute Ti through the re-solution and precipitation of Ti in the steel, and ensures good fatigue resistance after post-heat treatment. Is important to. When the extraction temperature is equal to or lower than the equilibrium solid solution temperature T Ti calculated from the following equation (1), coarse Ti precipitated during continuous casting remains as an undissolved carbonitride and becomes solid at the stage of the raw hot rolled steel sheet. The amount of dissolved Ti is less than 0.005%, and the fatigue resistance characteristics in the plastic strain region that are remarkably excellent after post-heat treatment cannot be obtained. By extracting the slab under a temperature condition higher than the equilibrium solid solution temperature T Ti calculated from the following equation (1), 0.005% or more of Ti exists as a solid solution Ti at the stage of the raw hot-rolled steel sheet. By the heat treatment after processing, 0.005% or more of Ti can be precipitated as fine precipitates having a particle diameter of 20 nm or less, and the fatigue resistance property in the plastic strain region that is remarkably excellent can be obtained. More preferably, the slab extraction temperature is preferably 1620 K or less from the viewpoint of preventing the crystal grain size from becoming coarse, and the slab leveling is ensured from the viewpoint of ensuring uniformity of the solid solution state of Ti and sufficient solid solution time. The heat time (time for holding the slab at a temperature higher than the equilibrium solid solution temperature T Ti ) is preferably 10 min or more.
ｌｏｇ（［Ｔｉ－Ｎ×４８÷１４］［Ｃ］）＝－７０００／（ＴＴｉ（℃）＋２７３）＋２．７５　・・・（１）
ただし、（１）式におけるＴｉ、Ｎ、Ｃは、鋼スラブ中のそれぞれの元素の含有量（質量％）である。 log ([Ti−N × 48 ÷ 14] [C]) = − 7000 / (T Ti (° C.) + 273) +2.75 (1)
However, Ti, N, and C in the formula (1) are contents (mass%) of respective elements in the steel slab.
　ＴＴｉ－４００℃以上の仕上げ圧延温度
　熱延仕上げ圧延温度がＴＴｉ－４００℃を下回ると、表面近傍部分での上下ロールによる付加的剪断歪、あるいはロールや冷却水による抜熱により歪誘起析出が誘発され、素材熱延鋼板の段階で特に表面近傍（表裏面から２００μｍ以内）に存在する固溶Ｔｉ量が０．００５％を下回り、後熱処理後に格段に優れた塑性歪域での耐疲労特性が得られない。熱延仕上げ圧延温度をＴＴｉ－４００℃以上とすることで、素材熱延鋼板の段階で表面近傍含め０．００５％以上のＴｉが固溶Ｔｉとして存在し、成形加工後の熱処理によって、０．００５％以上のＴｉを粒径２０ｎｍ以下の微細な析出物として析出させることができ、格段に優れた塑性歪域での耐疲労特性が得られる。 T Ti -400 ° C or higher finish rolling temperature When hot rolling finish rolling temperature is lower than T Ti -400 ° C, strain induced precipitation due to additional shear strain by upper and lower rolls near the surface or heat removal by rolls and cooling water. The amount of solid solution Ti present in the vicinity of the front surface (within 200 μm from the front and back surfaces) is less than 0.005% at the stage of the raw hot-rolled steel sheet, and fatigue resistance in a plastic strain region that is remarkably excellent after post-heat treatment Characteristics are not obtained. By setting the hot rolling finish rolling temperature to T Ti −400 ° C. or higher, 0.005% or more of Ti including the vicinity of the surface is present as solute Ti at the stage of the raw hot rolled steel sheet, and it is reduced to 0 by heat treatment after forming. 0.005% or more of Ti can be precipitated as fine precipitates having a particle size of 20 nm or less, and a particularly excellent fatigue resistance property in a plastic strain region can be obtained.
　ＴＴｉ－４００℃からＴＴｉ－５００℃までの温度域を１０℃／s以上の平均冷却速度で冷却
　ＴＴｉ－４００℃からＴＴｉ－５００℃までの温度域の平均冷却速度が１０℃／sを下回ると、ＴｉＣが熱延ランナウトからコイリングの過程で析出し、素材熱延鋼板の段階で存在する固溶Ｔｉ量が０．００５％を下回り、後熱処理後に格段に優れた塑性歪域での耐疲労特性が得られない。ＴＴｉ－４００℃からＴＴｉ－５００℃までの温度域を１０℃／ｓ以上の平均冷却速度で急冷することで、素材熱延鋼板の段階で表面近傍含め０．００５％以上のＴｉが固溶Ｔｉとして存在し、成形加工後の熱処理によって、０．００５％以上のＴｉを粒径２０ｎｍ以下の微細な析出物として析出させることができ、格段に優れた塑性歪域での耐疲労特性が得られる。 The average cooling rate of the temperature range of the temperature range from T Ti -400 ° C. until T Ti -500 ° C. from the cooling T Ti -400 ° C. at an average cooling rate of more than 10 ° C. / s until T Ti -500 ° C. is 10 ° C. / Below s, TiC precipitates from the hot-rolled runout during the coiling process, and the amount of solid solution Ti present at the stage of the raw hot-rolled steel sheet is less than 0.005%. Fatigue resistance characteristics cannot be obtained. By quenching the temperature range from T Ti −400 ° C. to T Ti −500 ° C. at an average cooling rate of 10 ° C./s or more, 0.005% or more of Ti including the vicinity of the surface is solidified at the stage of the hot rolled steel sheet. It exists as molten Ti, and 0.005% or more of Ti can be precipitated as fine precipitates with a particle size of 20 nm or less by heat treatment after forming, and has excellent fatigue resistance characteristics in the plastic strain region. can get.
　ＴＴｉ－５００℃以下の巻き取り温度
　巻き取り温度がＴＴｉ－５００℃を超えると、コイル冷却までの間にＴｉ析出物の析出が促進され、素材熱延鋼板の段階で存在する固溶Ｔｉ量が０．００５％を下回り、後熱処理後に格段に優れた塑性歪域での耐疲労特性が得られない。巻き取り温度をＴＴｉ－５００℃以下とすることで、素材熱延鋼板の段階で表面近傍含め０．００５％以上のＴｉが固溶Ｔｉとして存在し、成形加工後の熱処理によって、０．００５％以上のＴｉを粒径２０ｎｍ以下の微細な析出物として析出させることで、格段に優れた塑性歪域での耐疲労特性が得られる。なお、前記仕上げ圧延温度、巻き取り温度は、コイル幅中央部の表面温度であり、平均冷却速度は、前記表面温度から求められるものである。 Winding temperature of T Ti −500 ° C. or lower When the winding temperature exceeds T Ti −500 ° C., precipitation of Ti precipitates is promoted before coil cooling, and solid solution Ti present at the stage of the raw hot rolled steel sheet The amount is less than 0.005%, and the fatigue resistance property in the plastic strain region which is remarkably excellent after post heat treatment cannot be obtained. By setting the winding temperature to T Ti −500 ° C. or less, 0.005% or more of Ti including the vicinity of the surface is present as solute Ti at the stage of the raw hot-rolled steel sheet. By precipitating at least% Ti as fine precipitates having a particle size of 20 nm or less, excellent fatigue resistance characteristics in the plastic strain region can be obtained. The finish rolling temperature and the coiling temperature are surface temperatures at the center of the coil width, and the average cooling rate is obtained from the surface temperature.
　上記の製造方法により、組織中に０．００５％以上のＴｉが固溶Ｔｉとして存在する熱延鋼板（素材熱延鋼板）が得られる。 By the above manufacturing method, a hot-rolled steel sheet (raw material hot-rolled steel sheet) in which 0.005% or more of Ti exists as a solid solution Ti in the structure is obtained.
　次に、本発明の鋼部材の製造方法について説明する。
　本発明の鋼部材は、上記素材熱延鋼板に、成形加工を施した後、所定の熱処理を施すことで製造される。成形加工としては、特に限定されないが、例えば、鋼部材が鋼管であれば、造管加工が挙げられる。鋼部材が溶接鋼管であれば、造管加工後に溶接加工を施してもよい。また、例えば、鋼部材が自動車用構造部材等の成形部品であれば、プレス加工等が挙げられる。
　成形加工後、以下の条件で熱処理を施す。 Next, the manufacturing method of the steel member of this invention is demonstrated.
The steel member of the present invention is manufactured by subjecting the material hot-rolled steel sheet to a predetermined heat treatment after forming. Although it does not specifically limit as a shaping | molding process, For example, if a steel member is a steel pipe, a pipe making process will be mentioned. If the steel member is a welded steel pipe, the welding process may be performed after the pipe making process. In addition, for example, if the steel member is a molded part such as a structural member for an automobile, press working or the like can be given.
After forming, heat treatment is performed under the following conditions.
　５５０℃を超え１０５０℃以下の温度に加熱した後、５５０～４００℃の温度域を１０℃／ｓ以上の平均冷却速度で冷却
　素材熱延鋼板に成形加工を施した後、５５０℃を超え１０５０℃以下の温度に加熱した後、５５０～４００℃の温度域を１０℃／ｓ以上の平均冷却速度で冷却する熱処理を施すことで、０．００５％以上のＴｉが粒径２０ｎｍ以下の微細な析出物として析出し、格段に優れた塑性歪域での耐疲労特性が得られる。加熱温度が５５０℃以下であると固溶Ｔｉが２０ｎｍ以下の微細な析出物として析出せず、格段に優れた塑性歪域での耐疲労特性が得られない。また、加熱温度が１０５０℃を超えるとフェライト相の粒径が５０μｍを超え、格段に優れた塑性歪域での耐疲労特性が得られにくくなる。また５５０～４００℃の温度域の冷却速度が１０℃／ｓを下回ると十分な強度（ＹＳ≧７７０ＭＰａ）が得られない。なお、加熱温度はさらに望ましくは７００～１０００℃の範囲である。 After heating to a temperature exceeding 550 ° C. and not exceeding 1050 ° C., cooling the temperature range of 550 to 400 ° C. at an average cooling rate of 10 ° C./s or more. After forming the material hot-rolled steel sheet, exceeding 550 ° C. and exceeding 1050 After heating to a temperature of 550 ° C. or less, a heat treatment is performed to cool the temperature range of 550 to 400 ° C. at an average cooling rate of 10 ° C./s or more, whereby 0.005% or more of Ti has a fine particle size of 20 nm or less. Precipitates as precipitates and provides excellent resistance to fatigue in the plastic strain region. When the heating temperature is 550 ° C. or lower, the solid solution Ti is not precipitated as fine precipitates having a thickness of 20 nm or less, and the remarkably excellent fatigue resistance property in the plastic strain region cannot be obtained. On the other hand, when the heating temperature exceeds 1050 ° C., the particle size of the ferrite phase exceeds 50 μm, and it becomes difficult to obtain fatigue resistance characteristics in a plastic strain region that is remarkably excellent. If the cooling rate in the temperature range of 550 to 400 ° C. is less than 10 ° C./s, sufficient strength (YS ≧ 770 MPa) cannot be obtained. The heating temperature is more preferably in the range of 700 to 1000 ° C.
　また、特に限定されないが、例えば溶接鋼管を製造する場合には、素材熱延鋼板を、黒皮まま、或いは、必要に応じて、酸洗、冷間圧延、焼鈍、めっきのいずれかまたは複数の処理を行った後、スリッティングで所定の板幅とし、コイルを長手方向に１コイル以上溶接接合し、ロール成形或いはプレス成形により概円形断面成形に成形し、端部を高周波電縫溶接、レーザー溶接等の方法により接合し、オンライン或いはオフラインで５５０℃を超え１０５０℃以下の温度に加熱後、５５０～４００℃の温度域を１０℃／ｓ以上の平均冷却速度で冷却し、コイル状にまかれた鋼管とする。
　また、例えば成形部品を製造する場合には、素材熱延鋼板を、黒皮まま、或いは、必要に応じて、酸洗、冷間圧延、焼鈍、めっきのいずれかまたは複数の処理を行った後、所定の大きさにブランキングし、部品に成形加工した後、５５０℃を超え１０５０℃以下の温度に加熱後、５５０～４００℃の温度域を１０℃／ｓ以上の平均冷却速度で冷却する。これにより、０．００５％以上のＴｉが粒径２０ｎｍ以下の微細な析出物として析出し、格段に優れた塑性歪域での耐疲労特性が得られる。 Further, although not particularly limited, for example, when producing a welded steel pipe, the raw hot-rolled steel sheet is left as it is, or, if necessary, pickling, cold rolling, annealing, plating, or a plurality of After processing, slitting to a predetermined plate width, one or more coils are welded and joined in the longitudinal direction, formed into a roughly circular cross-section by roll forming or press forming, and the end is subjected to high-frequency electric seam welding, laser Joined by welding, etc., and heated online or offline to a temperature exceeding 550 ° C. and below 1050 ° C., cooling the temperature range of 550 to 400 ° C. at an average cooling rate of 10 ° C./s or more to form a coil. Steel pipe.
In addition, for example, when manufacturing a molded part, the raw hot-rolled steel sheet is left as it is, or after performing any one or more of pickling, cold rolling, annealing, plating as necessary. , Blanked to a predetermined size, molded into a part, heated to a temperature exceeding 550 ° C. and below 1050 ° C., and then cooled at a temperature range of 550 to 400 ° C. at an average cooling rate of 10 ° C./s or more. . As a result, 0.005% or more of Ti precipitates as fine precipitates having a particle size of 20 nm or less, and the fatigue resistance characteristics in the plastic strain region that are remarkably excellent are obtained.
（実施例１）
　表１に示す組成（鋼種Ｃ～Ｌ）の鋼スラブを、スラブ表面温度約１２２０℃、スラブ中心温度約１２１０℃で加熱炉より抽出し、仕上げ圧延圧下率：９１％、コイル幅中央部仕上げ圧延温度約８６０℃、コイル幅方向最低仕上げ圧延温度約８５０℃、ＴＴｉ－４００℃からＴＴｉ－５００℃までの温度域を約２０℃／ｓの平均冷却速度で冷却し、巻取り温度約５６０℃とする熱間圧延を行い素材熱延鋼板（板厚：約５ｍｍ、先後端部の板厚は長手中央部に対し約１０％厚い）とした（Ｎｏ．３～１２）。
　また、表１に示す組成（鋼種Ａ）の鋼スラブを、スラブ表面温度約１２５０℃、スラブ中心温度約１２４５℃で加熱炉より抽出したこと以外は、上記と同様にして素材熱延鋼板とし（Ｎｏ．１）、表１に示す組成（鋼種Ｂ、Ｍ）の鋼スラブを、スラブ表面温度約１３３５℃、スラブ中心温度約１３３５℃で加熱炉より抽出し、コイル幅中央部仕上げ圧延温度を約９４０℃としたこと以外は、上記と同様にして素材熱延鋼板とした（Ｎｏ．２、１３）。 Example 1
A steel slab having the composition shown in Table 1 (steel types C to L) was extracted from a heating furnace at a slab surface temperature of about 1220 ° C and a slab center temperature of about 1210 ° C, and finish rolling reduction: 91%, coil width center finish rolling temperature of about 860 ° C., coil width direction minimum finish rolling temperature of about 850 ° C., and cooled at an average cooling rate of T Ti -400 ° C. from T Ti -500 temperature range of about 20 ° C. / s up to ° C., the coiling temperature of about 560 A hot rolled steel sheet (sheet thickness: about 5 mm, the thickness of the front and rear end portions is about 10% thicker than the longitudinal center portion) was obtained by hot rolling at a temperature of 0 ° C. (No. 3 to 12).
Moreover, except having extracted the steel slab of the composition (steel type A) shown in Table 1 from the heating furnace at a slab surface temperature of about 1250 ° C. and a slab center temperature of about 1245 ° C., a raw hot-rolled steel sheet was obtained in the same manner as above ( No. 1) Steel slabs having the compositions shown in Table 1 (steel types B and M) were extracted from a heating furnace at a slab surface temperature of about 1335 ° C. and a slab center temperature of about 1335 ° C., and the coil width center finish rolling temperature was about Except having set it as 940 degreeC, it was set as the raw material hot-rolled steel plate similarly to the above (No. 2, 13).
　次いで、これら素材熱延鋼板に酸洗を施したのち、所定の幅寸法にスリット加工し、連続成形してオープン管とし、該オープン管を高周波抵抗溶接により電縫溶接して、幅絞り率４％で、外径φ５０．８ｍｍ肉厚約５ｍｍの溶接鋼管を得た。この溶接鋼管全体を連続的に高周波加熱し、加熱温度９２０℃、保持時間約５秒加熱した後、外面より水でミスト冷却を行い、５５０～４００℃の温度域を約５０℃／ｓの平均冷却速度で冷却する熱処理を施した。 Next, after pickling these hot-rolled steel sheets, slitting to a predetermined width, continuously forming an open pipe, the open pipe is electro-sealed by high-frequency resistance welding, and a width drawing ratio of 4 %, A welded steel pipe having an outer diameter of 50.8 mm and a thickness of about 5 mm was obtained. The whole welded steel pipe is continuously heated at high frequency, heated at a heating temperature of 920 ° C and a holding time of about 5 seconds, then cooled with water from the outer surface, and the temperature range of 550 to 400 ° C is an average of about 50 ° C / s. A heat treatment for cooling at a cooling rate was performed.
　これらの溶接鋼管から試験片を採取し、組織観察試験、析出物、固溶量の定量試験、引張試験、塑性歪域疲労試験、低温靱性試験を実施した。試験方法はつぎの通りとした。 Specimens were collected from these welded steel pipes and subjected to a structure observation test, a precipitate, a quantitative test of the solid solution amount, a tensile test, a plastic strain region fatigue test, and a low temperature toughness test. The test method was as follows.
（１）組織観察試験
　これら溶接鋼管の円周方向断面が観察面となるように組織観察試験片を採取して、研磨、ナイタール腐食して走査型電子顕微鏡（３０００倍）で組織を観察し、ＥＢＳＤ（Ｅｌｅｃｔｒｏｎ　ＢａｃｋＳｃａｔｔｅｒ　Ｄｉｆｆｒａｃｔｉｏｎ）法により隣接粒との傾角１５°以上を粒界としてフェライト相の平均粒径を求めた。なお、表面から板厚方向２００μｍまでの平均粒径として、板厚方向５０～２００μｍの間を５０μｍピッチで３点を測定し平均した値と、板厚中心部を中心に中心偏析部を避け、板厚方向に５０μｍピッチで３点を測定平均した値をそれぞれ求めた。 (1) Microstructure observation test Samples of microstructural observation specimens were collected so that the circumferential cross section of these welded steel pipes became the observation surface, polished, nital-corroded, and observed with a scanning electron microscope (3000 times). The average grain size of the ferrite phase was determined by an EBSD (Electron BackScatter Diffraction) method with an inclination angle of 15 ° or more with adjacent grains as a grain boundary. The average particle diameter from the surface to the plate thickness direction of 200 μm was measured by averaging three points at a pitch of 50 μm between the plate thickness directions of 50 to 200 μm and the center segregation portion around the plate thickness center, Values obtained by measuring and averaging three points at a pitch of 50 μm in the plate thickness direction were obtained.
（２）析出物、固溶量の定量試験
　これら溶接鋼管から、２０ｍｍ×３０ｍｍの大きさの試料片を切り出し、１０％ＡＡ系電解液（１０ｖｏｌ％アセチルアセトン－１ｍａｓｓ％塩化テトラメチルアンモニウム－メタノール）中で、約０．２ｇを電流密度２０ｍＡ／ｃｍ２で定電流電解した。電解後の、表面に析出物が付着している試料片を電解液から取り出して、ヘキサメタリン酸ナトリウム水溶液（５００ｍｇ／ｌ）（以下、ＳＨＭＰ水溶液と称す）中に浸漬し、超音波振動を付与して、析出物を試料片から剥離しＳＨＭＰ水溶液中に抽出した。次いで、析出物を含むＳＨＭＰ水溶液を、穴径１００ｎｍ、２０ｎｍの順にフィルタを用いてろ過し、ろ過後のフィルタ上の残渣とろ液に対してＩＣＰ発光分光分析装置を用いて分析し、フィルタ上の残渣中およびろ液中のＴｉの絶対量を測定し、粒径１００ｎｍを超える析出物、粒径１００ｎｍ以下２０ｎｍ超の析出物、粒径２０ｎｍ以下の析出物に含まれるＴｉの絶対量Tilp、Timp、Tispをそれぞれ得た。なお、電解質量は、析出物剥離後の試料片の質量を測定し、電解前の試料片の質量から差し引くことで求めた。 (2) Quantitative test of precipitates and solid solution amount From these welded steel pipes, a sample piece having a size of 20 mm × 30 mm was cut out and in a 10% AA electrolyte solution (10 vol% acetylacetone-1 mass% tetramethylammonium chloride-methanol). Then, about 0.2 g was subjected to constant current electrolysis at a current density of 20 mA / cm 2 . After the electrolysis, the sample piece with the deposit attached on the surface is taken out from the electrolytic solution and immersed in an aqueous solution of sodium hexametaphosphate (500 mg / l) (hereinafter referred to as an SHMP aqueous solution) to give ultrasonic vibration. The precipitate was peeled from the sample piece and extracted into an aqueous SHMP solution. Next, the SHMP aqueous solution containing the precipitate is filtered using a filter in the order of hole diameters of 100 nm and 20 nm, and the residue and filtrate on the filtered filter are analyzed using an ICP emission spectroscopic analyzer. The absolute amount of Ti in the residue and the filtrate is measured, and the absolute amount of Ti contained in precipitates having a particle size of more than 100 nm, precipitates having a particle size of 100 nm or less and more than 20 nm, and precipitates having a particle size of 20 nm or less. , Tisp was obtained respectively. In addition, the electrolytic mass was calculated | required by measuring the mass of the sample piece after deposit peeling, and subtracting from the mass of the sample piece before electrolysis.
　固溶状態にあるＴｉ（固溶Ｔｉ）は、電解後の電解液を分析溶液とし、ＩＣＰ質量分析法を用いてＴｉ及び比較元素としてＦｅの液中濃度を測定した。得られた濃度を基に、Ｆｅに対するＴｉの濃度比をそれぞれ算出し、さらに、試料中のＦｅの含有率を乗じることで、固溶状態にあるＴｉの含有率を求めた。なお、試料中のＦｅの含有率は、Ｆｅ以外の組成値の合計を１００％から減算することで求めることができる。この析出物、固溶量の定量試験は、後熱処理を施した後の溶接鋼管に加えて、後熱処理を施す前の溶接鋼管についても行った。 For Ti (solid solution Ti) in a solid solution state, the electrolytic solution after electrolysis was used as an analysis solution, and the concentration of Ti and Ti as a comparative element in the solution was measured using ICP mass spectrometry. Based on the obtained concentration, the concentration ratio of Ti to Fe was calculated, and the content ratio of Ti in a solid solution state was determined by multiplying the content ratio of Fe in the sample. In addition, the content rate of Fe in a sample can be calculated | required by subtracting the sum total of composition values other than Fe from 100%. In addition to the welded steel pipe after the post-heat treatment, the quantitative test of the precipitate and the solid solution amount was also performed on the welded steel pipe before the post-heat treatment.
（３）引張試験
　これら溶接鋼管から、Ｌ方向が引張方向となるように、ＪＩＳ　Ｚ　２２０１の規定に準拠してＪＩＳ　１２号試験片を切出し、ＪＩＳ　Ｚ　２２４１の規定に準拠して引張試験を実施し、引張特性（引張強さＴＳ、降伏強さＹＳ、全伸びＥｌ）を求めた。 (3) Tensile test From these welded steel pipes, a JIS No. 12 test piece is cut out in accordance with JIS Z 2201 so that the L direction is the tensile direction, and a tensile test is performed in accordance with JIS Z 2241. The tensile properties (tensile strength TS, yield strength YS, total elongation El) were determined.
（４）塑性歪域疲労試験
　これら溶接鋼管から、板厚約５ｍｍ×板幅５ｍｍ、平行部長さ１２ｍｍの平行部断面寸法の板状Ｌ方向疲労試験片を偏平矯正後に採取し、引張モード、ひずみ制御モード、ひずみ比＝０、全ひずみ範囲２．０％、サイクル数０．１２５Ｈｚの条件で疲労試験を行った。引張最大荷重が初期荷重から２５％低下したサイクル数を求め、破断までの繰り返し数とした。 (4) Plastic strain region fatigue test From these welded steel pipes, a plate-shaped L direction fatigue test piece having a plate section of about 5 mm × plate width 5 mm and a parallel section length of 12 mm was sampled after flattening, tensile mode, strain A fatigue test was performed under the conditions of control mode, strain ratio = 0, total strain range 2.0%, cycle number 0.125 Hz. The number of cycles in which the maximum tensile load was reduced by 25% from the initial load was determined and used as the number of repetitions until breakage.
（５）低温靭性試験
　これら溶接鋼管から管長手方向（Ｌ方向）が試験片長さとなるように展開し、ＪＩＳ　Ｚ　２２０２の規定に準拠してシャルピー試験片（２ｍｍＶノッチ、１／２サイズ）を切出し、ＪＩＳ　Ｚ　２２４２の規定に準拠してシャルピー衝撃試験を実施し、破面遷移温度を求め、低温靭性を評価した。 (5) Low temperature toughness test Expanded from these welded steel pipes so that the longitudinal direction of the pipe (L direction) is the length of the specimen, and cut out Charpy specimen (2mmV notch, 1/2 size) in accordance with JIS Z 2202 regulations. The Charpy impact test was performed in accordance with the provisions of JIS Z 2242, the fracture surface transition temperature was determined, and the low temperature toughness was evaluated.
　また、上述の方法により、表面から板厚方向２００μｍまでの平均硬度（HV(0.1)S）、中心偏析部を除く板厚中心近傍の平均硬度（HV(0.1)C）を測定し、表面から板厚方向２００μｍまでの平均硬度と、中心偏析部を除く板厚中心近傍の平均硬度の差ΔＨＶ（HV(0.1)C－HV(0.1)S）を求めた。
　得られた結果を表２に示す。 In addition, the average hardness (HV (0.1) S ) from the surface to the thickness direction of 200 μm and the average hardness near the center of the plate thickness (HV (0.1) C ) excluding the central segregation part are measured by the above method. A difference ΔHV (HV (0.1) C −HV (0.1) S ) between the average hardness in the thickness direction of 200 μm and the average hardness in the vicinity of the thickness center excluding the center segregation portion was determined.
The obtained results are shown in Table 2.
Figure JPOXMLDOC01-appb-T000001
  
Figure JPOXMLDOC01-appb-T000001
  
Figure JPOXMLDOC01-appb-T000002
  
Figure JPOXMLDOC01-appb-T000002
  
　本発明例（Ｎｏ．１～１１）は、いずれも上述の塑性歪域疲労試験におけるサイクル数が１０００サイクル以上となり、塑性歪域での耐疲労特性に優れる。さらに、本発明例は、いずれもＹＳが７７０ＭＰａ以上であり強度特性にも優れる。また、本発明例は、いずれもシャルピー破面遷移温度が－３０℃以下であり低温靭性にも優れる。一方、鋼の成分組成が本発明の範囲を満たさず粒径２０ｎｍ以下の析出物として析出したＴｉが０．００５％未満であるＮｏ．１２、鋼の成分組成が本発明の範囲を満たさないＮｏ．１３は、所望の塑性歪域での耐疲労特性が得られない。 In each of the inventive examples (Nos. 1 to 11), the number of cycles in the above-described plastic strain region fatigue test is 1000 cycles or more, and the fatigue resistance property in the plastic strain region is excellent. Further, in all of the examples of the present invention, YS is 770 MPa or more and excellent in strength characteristics. Further, in all of the examples of the present invention, the Charpy fracture surface transition temperature is −30 ° C. or lower, and the low temperature toughness is excellent. On the other hand, the composition of steel does not satisfy the scope of the present invention, and Ti deposited as a precipitate having a particle size of 20 nm or less is less than 0.005%. No. 12, steel component composition does not meet the scope of the present invention. No. 13 does not provide fatigue resistance characteristics in the desired plastic strain region.
（実施例２）
　表１に示す鋼種Ａ、Ｂ、Ｃの成分組成を有する鋼スラブに、表３に示す条件の熱間圧延を施し素材熱延鋼板（板厚：約５ｍｍ、先後端部の板厚は長手中央部に対し約１０％厚い）とした。ついでこれら素材熱延鋼板に酸洗を施したのち、所定の幅寸法にスリット加工し、連続成形してオープン管とし、該オープン管を高周波抵抗溶接により電縫溶接して、幅絞り率４％で、外径φ５０．８ｍｍ肉厚約５ｍｍの溶接鋼管を得た。この溶接鋼管全体を連続的に高周波加熱し、表３に示す条件で熱処理を施した。これらの溶接鋼管から試験片を採取し、組織観察試験、析出物、固溶量の定量試験、引張試験、塑性歪域疲労試験、低温靱性試験、ビッカース硬さ測定を実施した。
　なお、Ｎｏ．２３については、素材熱延鋼板に酸洗を施したのち、所定の大きさにブランキングし、プレス加工して成形部品としたものに対して、表３に示す条件で熱処理を施した。そして、この成形部品から試験片を採取し、上記各試験を実施した。 (Example 2)
A steel slab having the composition of steel types A, B, and C shown in Table 1 is subjected to hot rolling under the conditions shown in Table 3 and a hot-rolled steel sheet (thickness: about 5 mm, the thickness at the front and rear ends is the longitudinal center) About 10% thicker). Next, after pickling these hot-rolled steel sheets, slitting them to a predetermined width, continuously forming an open pipe, the open pipe is electro-welded by high-frequency resistance welding, and the width drawing ratio is 4%. Thus, a welded steel pipe having an outer diameter of 50.8 mm and a thickness of about 5 mm was obtained. The entire welded steel pipe was continuously heated at a high frequency and heat treated under the conditions shown in Table 3. Specimens were collected from these welded steel pipes, and subjected to a structure observation test, a precipitate, a quantitative test of the solid solution amount, a tensile test, a plastic strain region fatigue test, a low temperature toughness test, and a Vickers hardness measurement.
In addition, No. For No. 23, the hot-rolled steel sheet was pickled, blanked to a predetermined size, and subjected to a heat treatment under the conditions shown in Table 3 for a molded part by pressing. And the test piece was extract | collected from this molded part, and each said test was implemented.
　得られた結果を表４に示す。なお、表３、表４には、上記Ｎｏ．１～３の結果を併記してある。 Table 4 shows the obtained results. In Tables 3 and 4, the above No. Results 1 to 3 are also shown.
Figure JPOXMLDOC01-appb-T000003
  
Figure JPOXMLDOC01-appb-T000003
  
Figure JPOXMLDOC01-appb-T000004
  
Figure JPOXMLDOC01-appb-T000004
  
　本発明例（Ｎｏ．２１～２３）はいずれも、上述の塑性歪域疲労試験におけるサイクル数が１０００サイクル以上となり、塑性歪域での耐疲労特性に優れる。さらに、本発明例は、いずれもＹＳが７７０ＭＰａ以上であり強度特性にも優れる。また、本発明例は、いずれもシャルピー破面遷移温度が－３０℃以下であり低温靭性にも優れる。一方、粒径２０ｎｍ以下の析出物として析出したＴｉ量が本発明の範囲外であるＮｏ．１４～２０は、所望の塑性歪域での耐疲特性が得られない。 In all of the inventive examples (Nos. 21 to 23), the number of cycles in the plastic strain region fatigue test is 1000 cycles or more, and the fatigue resistance property in the plastic strain region is excellent. Further, in all of the examples of the present invention, YS is 770 MPa or more and excellent in strength characteristics. Further, in all of the examples of the present invention, the Charpy fracture surface transition temperature is −30 ° C. or lower, and the low temperature toughness is excellent. On the other hand, the amount of Ti deposited as a precipitate having a particle size of 20 nm or less is outside the scope of the present invention. In Nos. 14 to 20, fatigue resistance characteristics in a desired plastic strain region cannot be obtained.

Claims (9)

  1. 　質量％で、Ｔｉを０．０３１～０．２００％含有し、組織中に０．００５％以上のＴｉが粒径２０ｎｍ以下の析出物として析出している、鋼部材。 A steel member containing 0.031 to 0.200% Ti by mass%, and 0.005% or more of Ti is precipitated as a precipitate having a particle size of 20 nm or less in the structure.
  2. 　前記鋼部材は、質量％で、
    Ｃ：０．１９～０．５０％、
    Ｓｉ：０．００２～１．５％、
    Ｍｎ：０．４～２．５％、
    Ａｌ：０．０１～０．１９％、
    Ｃｒ：０．００１～０．９０％、
    Ｂ：０．０００１～０．００５０％、
    Ｔｉ：０．０３１～０．２００％、
    Ｐ：０．０１９％以下（０％を含む）、
    Ｓ：０．０１５％以下（０％を含む）、
    Ｎ：０．００８％以下（０％を含む）、
    Ｏ：０．００３％以下（０％を含む）、
    Ｓｎ：０．１０％以下（０％を含む）を含有し、
    残部がＦｅおよび不可避的不純物からなる組成を有する、請求項１に記載の鋼部材。     The steel member is in mass%,
    C: 0.19 to 0.50%,
    Si: 0.002 to 1.5%,
    Mn: 0.4 to 2.5%,
    Al: 0.01 to 0.19%,
    Cr: 0.001 to 0.90%,
    B: 0.0001 to 0.0050%,
    Ti: 0.031 to 0.200%,
    P: 0.019% or less (including 0%),
    S: 0.015% or less (including 0%),
    N: 0.008% or less (including 0%),
    O: 0.003% or less (including 0%),
    Sn: 0.10% or less (including 0%),
    The steel member according to claim 1, wherein the balance has a composition comprising Fe and inevitable impurities.
  3. 　前記組成に加えてさらに、質量％で、
    Ｎｂ：０．００１～０．１５％、
    Ｖ：０．００１～０．１５％、
    Ｗ：０．００１～０．１５％、
    Ｍｏ：０．００１～０．４５％、
    Ｃｕ：０．００１～０．４５％、
    Ｎｉ：０．００１～０．４５％、
    Ｃａ：０．０００１～０．００５％、
    Ｓｂ：０．０００１～０．１０％
    のうちから選ばれた１種または２種以上を含有する、請求項２に記載の鋼部材。     In addition to the above composition,
    Nb: 0.001 to 0.15%,
    V: 0.001 to 0.15%,
    W: 0.001 to 0.15%,
    Mo: 0.001 to 0.45%,
    Cu: 0.001 to 0.45%,
    Ni: 0.001 to 0.45%,
    Ca: 0.0001 to 0.005%,
    Sb: 0.0001 to 0.10%
    The steel member according to claim 2, comprising one or more selected from among the above.
  4. 　前記鋼部材が溶接鋼管である、請求項１～３のいずれかに記載の鋼部材。 The steel member according to any one of claims 1 to 3, wherein the steel member is a welded steel pipe.
  5. 　請求項１～４のいずれかに記載の鋼部材用の熱延鋼板であって、
    　質量％で、Ｔｉを０．０３１～０．２００％含有し、組織中に０．００５％以上のＴｉが固溶Ｔｉとして存在する、鋼部材用の熱延鋼板。     A hot-rolled steel sheet for a steel member according to any one of claims 1 to 4,
    A hot-rolled steel sheet for steel members, containing 0.031 to 0.200% Ti by mass and 0.005% or more of Ti present as a solid solution Ti in the structure.
  6. 　長手方向両端部である先端部および尾端部の板厚が、ともに長手方向中央部の板厚に比べて５～５０％厚い、請求項５に記載の鋼部材用の熱延鋼板。 6. The hot-rolled steel sheet for steel members according to claim 5, wherein the thickness of the tip and tail ends, which are both ends in the longitudinal direction, is 5 to 50% thicker than the thickness of the central portion in the longitudinal direction.
  7. 　請求項１～４のいずれかに記載の鋼部材の製造方法であって、
    　質量％で、Ｔｉを０．０３１～０．２００％含有し、組織中に０．００５％以上のＴｉが固溶Ｔｉとして存在する熱延鋼板に成形加工を施した後に、５５０℃を超え１０５０℃以下の温度に加熱した後、５５０～４００℃の温度域を１０℃／ｓ以上の平均冷却速度で冷却する熱処理を施す、鋼部材の製造方法。     A method for producing a steel member according to any one of claims 1 to 4,
    After forming a hot-rolled steel sheet containing 0.031 to 0.200% Ti by mass and 0.005% or more of Ti as a solid solution Ti in the structure, the temperature exceeds 550 ° C. and exceeds 1050 A method for producing a steel member, wherein the steel member is heated to a temperature of ℃ or less and then subjected to a heat treatment for cooling a temperature range of 550 to 400 ℃ at an average cooling rate of 10 ℃ / s or more.
  8. 　前記熱延鋼板を、質量％で、Ｔｉを０．０３１～０．２００％含有する鋼スラブを下記（１）式から計算される平衡固溶温度ＴＴｉよりも高い温度条件でスラブ抽出した後、ＴＴｉ－４００℃以上の温度で仕上げ圧延を終了し、ＴＴｉ－４００℃からＴＴｉ－５００℃までの温度域を１０℃／ｓ以上の平均冷却速度で冷却し、ＴＴｉ－５００℃以下の温度で巻き取って製造する、請求項７に記載の鋼部材の製造方法。
    ｌｏｇ（［Ｔｉ－Ｎ×４８÷１４］［Ｃ］）＝－７０００／（ＴＴｉ（℃）＋２７３）＋２．７５　・・・（１）
    ただし、（１）式におけるＴｉ、Ｎ、Ｃは、鋼スラブ中のそれぞれの元素の含有量（質量％）である。     After the hot-rolled steel sheet is slab-extracted at a temperature condition higher than the equilibrium solid solution temperature T Ti calculated from the following formula (1), a steel slab containing 0.031 to 0.200% Ti by mass% ends the finish rolling at T Ti -400 ° C. or higher, the temperature range from T Ti -400 ° C. until T Ti -500 ° C. and cooled at 10 ° C. / s or more average cooling rate, T Ti -500 ° C. The manufacturing method of the steel member of Claim 7 which winds and manufactures at the following temperatures.
    log ([Ti−N × 48 ÷ 14] [C]) = − 7000 / (T Ti (° C.) + 273) +2.75 (1)
    However, Ti, N, and C in the formula (1) are contents (mass%) of respective elements in the steel slab.
  9. 　請求項５または６に記載の鋼部材用の熱延鋼板の製造方法であって、
    　質量％で、Ｔｉを０．０３１～０．２００％含有する鋼スラブを、下記（１）式から計算される平衡固溶温度ＴＴｉよりも高い温度条件でスラブ抽出した後、ＴＴｉ－４００℃以上の温度で仕上げ圧延を終了し、ＴＴｉ－４００℃からＴＴｉ－５００℃までの温度域を１０℃／ｓ以上の平均冷却速度で冷却し、ＴＴｉ－５００℃以下の温度で巻き取る、鋼部材用の熱延鋼板の製造方法。
    ｌｏｇ（［Ｔｉ－Ｎ×４８÷１４］［Ｃ］）＝－７０００／（ＴＴｉ（℃）＋２７３）＋２．７５　・・・（１）
    ただし、（１）式におけるＴｉ、Ｎ、Ｃは、鋼スラブ中のそれぞれの元素の含有量（質量％）である。     A method for producing a hot-rolled steel sheet for steel members according to claim 5 or 6,
    A steel slab containing 0.031 to 0.200% Ti by mass is slab-extracted under a temperature condition higher than the equilibrium solid solution temperature T Ti calculated from the following equation (1), and then T Ti -400 Finish rolling at a temperature not lower than ° C. and finish cooling at a temperature range from T Ti −400 ° C. to T Ti −500 ° C. at an average cooling rate of 10 ° C./s or higher and winding at a temperature not higher than T Ti −500 ° C. A method for producing a hot-rolled steel sheet for a steel member.
    log ([Ti−N × 48 ÷ 14] [C]) = − 7000 / (T Ti (° C.) + 273) +2.75 (1)
    However, Ti, N, and C in the formula (1) are contents (mass%) of respective elements in the steel slab.
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