CN110494582B - Steel member, hot-rolled steel sheet for said steel member, and method for producing same - Google Patents

Steel member, hot-rolled steel sheet for said steel member, and method for producing same Download PDF

Info

Publication number
CN110494582B
CN110494582B CN201880023595.9A CN201880023595A CN110494582B CN 110494582 B CN110494582 B CN 110494582B CN 201880023595 A CN201880023595 A CN 201880023595A CN 110494582 B CN110494582 B CN 110494582B
Authority
CN
China
Prior art keywords
temperature
hot
steel
steel member
steel sheet
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Active
Application number
CN201880023595.9A
Other languages
Chinese (zh)
Other versions
CN110494582A (en
Inventor
丰田俊介
杉本一郎
川村修司
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
JFE Steel Corp
Original Assignee
JFE Steel Corp
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by JFE Steel Corp filed Critical JFE Steel Corp
Publication of CN110494582A publication Critical patent/CN110494582A/en
Application granted granted Critical
Publication of CN110494582B publication Critical patent/CN110494582B/en
Active legal-status Critical Current
Anticipated expiration legal-status Critical

Links

Images

Classifications

    • 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

Abstract

Provided are a steel member having excellent fatigue resistance in a plastic strain region, a hot-rolled steel sheet as a material for the steel member, and methods for producing the steel member and the hot-rolled steel sheet. A steel member containing 0.031-0.200% by mass of Ti, wherein 0.005% or more of Ti is precipitated in the form of precipitates having a grain size of 20nm or less. The hot-rolled steel sheet for steel members contains 0.031-0.200% by mass of Ti, and 0.005% or more of Ti is present in the structure as solid-solution Ti. In the method for manufacturing a steel member, the hot-rolled steel sheet is subjected to forming and then heat treatment as follows: heating to a temperature higher than 550 ℃ and lower than 1050 ℃, and then cooling at an average cooling rate of 10 ℃/sec or higher within a temperature range of 550-400 ℃. The method for producing the hot-rolled steel sheet comprises subjecting a steel slab containing 0.031-0.200% by mass of Ti to a temperature higher than the equilibrium solid solution temperature T determined by a predetermined formulaTiAfter billet extraction at the temperature of (1), at TTiFinishing the finish rolling at a temperature of-400 ℃ or higher, at a temperature of from TTi-400 ℃ to TTiCooling at a temperature in the range of-500 ℃ at an average cooling rate of 10 ℃/s or more, at TTiCoiling is carried out at a temperature below-500 ℃.

Description

Steel member, hot-rolled steel sheet for said steel member, and method for producing same
Technical Field
The present invention relates to a steel member, a hot-rolled steel sheet for the steel member, and methods for producing the same. More specifically, the present invention relates to a steel member excellent in fatigue resistance in a plastic strain region, a hot-rolled steel sheet for the steel member, and methods for producing the same. The present invention particularly relates to a welded steel pipe for a coiled tubing, a welded steel pipe for a line pipe, and a welded steel pipe for a structural member for an automobile, which require high strength and fatigue resistance in a plastic strain region, and particularly relates to a welded steel pipe for a coiled tubing, and to an improvement in fatigue life of these steel members in a plastic strain region.
Background
Patent document 1 discloses a method for producing a high-tension electric resistance welded steel pipe having a strength after pipe production of 700MPa or more in yield strength and 800MPa or more in tensile strength and ductility of 15% or more in elongation, which is used as a high-strength structural member and a driving force transmission member of an automobile or the like, or as an electric resistance welded pipe for cleaning an oil country tubular good. According to this method, a high-tension electric resistance welded steel pipe that does not cause softening of the weld heat affected zone can be obtained by containing 0.09 to 0.18% of C and predetermined amounts of the alloying elements Cu, Ni, Cr, and Mo. However, as a steel pipe for a coiled tubing which is used for fatigue applications, particularly, which requires fatigue resistance in a plastic strain region, there is a problem that the durability life is low when repeatedly used.
Patent document 2 discloses a steel strip for a continuous oil pipe excellent in material uniformity and a method for producing the same. According to this method, a steel strip for a coiled tubing having small variations in yield strength in the coil width direction and the coil length direction can be obtained by containing 0.10 to 0.16% of C and predetermined amounts of the alloying elements Cr, Cu, Ni, Mo, Nb, and Ti. However, the fatigue resistance in the plastic strain region is insufficient, and the durability life in repeated use is low.
Patent document 3 discloses a quenched-tempered steel pipe excellent in fatigue life as a steel pipe for a mechanical structure of an automobile or the like, particularly for an automobile hollow stabilizer. According to this method, a steel pipe having a high fatigue life can be obtained by containing a predetermined chemical component, making the average grain size of precipitated carbides 0.5 μm or less, and making the hardness of the central portion of the wall thickness 400 HV. However, the level of fatigue life obtained for this steel pipe is the elastic region fatigue characteristic of low stress-high cycle with a lifetime of several tens of thousands of cycles. On the other hand, coiled tubing is repeatedly inserted into and recovered from a well and used several hundred times. A strain of about 2% plastic region is applied to the unwinding-winding of the coil and the bending (gooseneck) portion when inserting into the well, requiring high strain-low cycle fatigue strength of 100-1000 cycles. In general, the fatigue strength under a condition where the stress amplitude is constant, such as the elastic region fatigue, is increased by increasing the material strength. On the other hand, the strain in the longitudinal direction applied to the coiled tubing corresponds to a strain constant condition determined by the inner diameters of the coil and the gooseneck, and the contribution of the fatigue ductility coefficient of the so-called Morrow formula increases, so that there is a problem that the increase in strength does not necessarily lead to an improvement in the life, and the fatigue resistance characteristic in the desired plastic strain region cannot be obtained.
Documents of the prior art
Patent document
Patent document 1: japanese patent No. 3491339
Patent document 2: japanese patent No. 5494895
Patent document 3: japanese patent No. 5196934
Disclosure of Invention
Problems to be solved by the invention
The purpose 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 material for the steel member, and methods for producing the steel member and the hot-rolled steel sheet.
In the present invention, "excellent fatigue resistance in a plastic strain region" or "excellent fatigue resistance in a plastic strain region" means that the number of repetitions until fracture is 1000 or more when a tensile fatigue test is performed under the conditions of a tensile mode, a strain control mode, a strain ratio of 0, and a total strain range of 2.0%.
In addition, a hot-rolled steel sheet that is a material of the steel member of the present invention is also referred to as a "raw material hot-rolled steel sheet".
Examples of the steel member of the present invention include a steel pipe such as a welded steel pipe, a formed part such as an automotive structural member, and the like. Examples of the welded steel pipe include a welded steel pipe for a coiled tubing, a welded steel pipe for a line pipe, and a welded steel pipe for an automotive structural member.
Means for solving the problems
The present inventors have conducted systematic experimental studies by variously changing the chemical composition and the production conditions of a hot-rolled steel sheet as a material in order to achieve both of opposite properties of strength and fatigue resistance in a plastic strain region at a high level. As a result, they have found that a steel member satisfying both high strength and excellent fatigue resistance in a plastic strain region can be obtained by subjecting a steel having a specific chemical composition to hot rolling under specific temperature working conditions, forming into a steel pipe shape, or the like, and then heat treating under specific conditions.
The present invention has been completed based on such findings, and has the following configurations [1] to [9 ].
[1] A steel member containing 0.031-0.200% by mass of Ti, wherein 0.005% or more of Ti is precipitated in the form of precipitates having a grain size of 20nm or less.
[2] The steel member according to [1], wherein the above steel member has a composition as follows:
contains, in mass%, C: 0.19 to 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 to 0.0050%, Ti: 0.031-0.200%, P: less than 0.019% (including 0%), S: 0.015% or less (including 0%), N: less than 0.008% (including 0%), O: 0.003% or less (including 0%), Sn: 0.10% or less (including 0%), and the balance being Fe and unavoidable impurities.
[3] The steel member according to [2], further comprising an additive selected from the group consisting of Nb: 0.001-0.15%, V: 0.001-0.15%, W: 0.001 to 0.15%, Mo: 0.001 to 0.45%, Cu: 0.001-0.45%, Ni: 0.001-0.45%, Ca: 0.0001 to 0.005%, Sb: 0.0001-0.10% of one or more than two.
[4] The steel member according to any one of [1] to [3], wherein the steel member is a welded steel pipe.
[5] A hot-rolled steel sheet for a steel member according to any one of the above items [1] to [4], wherein 0.031 to 0.200% by mass of Ti is contained, and 0.005% or more of Ti is present in the structure as solid-solution Ti.
[6] A hot-rolled steel sheet for a steel member as recited in item [5], wherein the thickness of each of the front end portion and the rear end portion, which are both longitudinal end portions, is 5 to 50% thicker than the thickness of the central portion in the longitudinal direction.
[7] A method for producing a steel member according to any one of the above [1] to [4], wherein a hot-rolled steel sheet containing 0.031 to 0.200% by mass of Ti and 0.005% or more of Ti present in a structure as solid solution Ti is subjected to forming, and then subjected to heat treatment as follows: heating to a temperature higher than 550 ℃ and lower than 1050 ℃, and then cooling at an average cooling rate of 10 ℃/sec or higher within a temperature range of 550-400 ℃.
[8]Such as [7]]The method for manufacturing a steel member comprises subjecting a steel slab containing 0.031-0.200% by mass of Ti to a temperature higher than the equilibrium solid solution temperature T calculated by the following formula (1)TiAfter billet extraction at the temperature of (1), at TTiFinishing the finish rolling at a temperature of-400 ℃ or higher, at a temperature of from TTi-400 ℃ to TTiCooling at a temperature in the range of-500 ℃ at an average cooling rate of 10 ℃/s or more, at TTiAnd coiling the steel sheet at a temperature of-500 ℃ or lower to produce the hot-rolled steel sheet.
log([Ti-N×48÷14][C])＝-7000/(TTi(℃)+273)+2.75…(1)
Wherein Ti, N and C in the formula (1) represent the content (mass%) of each element in the steel slab.
[9]A method for producing a hot-rolled steel sheet for a steel member, which comprises the step of [5]]Or [6]]The method for producing a hot-rolled steel sheet for a steel member comprises subjecting a steel slab containing 0.031-0.200% by mass of Ti to a temperature higher than the equilibrium solid solution temperature T calculated by the following formula (1)TiAfter billet extraction at the temperature of (1), at TTiFinishing the finish rolling at a temperature of-400 ℃ or higher, at a temperature of from TTi-400 ℃ to TTiCooling at a temperature in the range of-500 ℃ at an average cooling rate of 10 ℃/s or more, at TTiCoiling is carried out at a temperature below-500 ℃.
log([Ti-N×48÷14][C])＝-7000/(TTi(℃)+273)+2.75…(1)
Wherein Ti, N and C in the formula (1) represent the content (mass%) of each element in the steel slab.
Effects of the invention
According to the present invention, a steel member having excellent fatigue resistance in a plastic strain region can be provided. The hot-rolled steel sheet of the present invention is particularly suitable as a material for the steel member.
According to the present invention, a steel member having high levels of both strength and fatigue resistance in a plastic strain region can be provided. Therefore, as the steel member of the present invention, a welded steel pipe for a coiled tubing, a welded steel pipe for a line pipe, and a welded steel pipe for a structural member for an automobile, which require high strength and fatigue resistance in a plastic strain region, are particularly preferable, and among them, a welded steel pipe for a coiled tubing is preferable.
Drawings
Fig. 1 is a graph showing the relationship between the amount of Ti precipitated as precipitates having a grain size of 20nm or less by the post heat treatment and the fatigue characteristics in the plastic strain region.
Detailed Description
(Steel member)
The steel member of the present invention is obtained by subjecting a hot-rolled steel sheet (raw hot-rolled steel sheet) produced by hot rolling under specific temperature processing conditions to forming and then heat treating under specific conditions. Hereinafter, the heat treatment performed after the forming process is performed on the raw hot-rolled steel sheet is also referred to as "post heat treatment".
First, the reason for limiting the chemical composition range of the steel member of the present invention will be described. In the following, the composition is expressed in% by mass.
Ti：0.031～0.200％
Ti precipitates as carbonitride in the hot rolling step, and suppresses recovery and recrystallization grain growth in the hot rolling step. The inclusion of Ti has an effect of obtaining a desired fine ferrite phase particle diameter (1 to 50 μm) in the structure (microstructure) of the raw hot-rolled steel sheet. The refinement of the microstructure at the stage of the raw hot-rolled steel sheet leads to the refinement of the microstructure after heat treatment after forming (cold working) such as pipe making and part forming, and excellent fatigue resistance in the plastic strain region can be obtained.
Tanaka et al propose the following model: dislocations are irreversibly accumulated on the sliding surface by fatigue cycles, and when the stress generated at this time exceeds the ultimate stress, initial cracks are generated (document: K. tanaka and T. Mura: J Appl Mech., Vol.48, p.97-103 (1981)). According to this model, G: transverse elastic constant, Ws: destruction energy per unit area, ν: poisson ratio, Δ τ: decomposed shear stress range on sliding surface, k: when the material physical property values such as the frictional force of dislocations on the sliding surface, the external force conditions, and the like are the same, the shorter the sliding surface length d, that is, the smaller the crystal grain size, the longer the fatigue crack generation period Nc of each crystal grain.
From such a mechanism, it is considered that the micro-structure material after the miniaturization of the present invention has a slow fatigue crack generation and exhibits excellent fatigue resistance in a plastic strain region.
Further, Ti is an essential element as described below: the strength after heat treatment after forming such as pipe making or component forming is improved by precipitation strengthening of the matrix as carbides, solid solution strengthening as solid solution elements, and strengthening of the phase transformation structure as hardenability-improving elements, and the fatigue strength is remarkably improved. Such an effect can be 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, at the stage of the raw material hot-rolled steel sheet described later, 0.005% or more of Ti exists as solid-solution Ti, and 0.005% or more of Ti cannot be precipitated as fine precipitates having a particle size of 20nm or less by heat treatment after forming, and the above effect cannot be obtained. On the other hand, when the Ti content exceeds the upper limit of the above range, coarse TiN is formed, and the fatigue resistance is lowered. Therefore, the Ti content is set to be in the range of 0.031-0.200%. The Ti content is preferably more than 0.120%. The Ti content is preferably 0.150% or less.
In the structure of the steel member of the present invention, 0.005% or more of Ti is precipitated as precipitates having a grain size of 20nm or less.
The present inventors have found that, when fatigue resistance in the plastic strain region is required after heat treatment (post heat treatment) performed after forming such as pipe making or part forming using a hot-rolled steel sheet as a material as in the present invention, 0.005% or more of Ti is precipitated as fine precipitates having a particle size of 20nm or less by the post heat treatment, whereby fatigue resistance in the plastic strain region which is particularly excellent can be obtained. Fig. 1 shows the relationship between the amount (mass%) of Ti precipitated as fine precipitates having a particle size of 20nm or less by the post heat treatment and the fatigue characteristics in the plastic strain region. When the amount of Ti precipitated as fine precipitates having a particle size of 20nm or less by the post heat treatment is 0.005% or more, the number of repetitions until fracture in the tensile fatigue test under the conditions of 0 strain ratio in the tensile mode, the strain control mode, and the total strain range of 2.0% is 1000 or more, and excellent fatigue resistance in the plastic strain region can be obtained.
Next, a preferred composition of the steel member of the present invention will be described.
C：0.19～0.50％
In the present invention, C is an element described as follows: by performing the post-heat treatment under specific conditions, high strength is ensured, and further, fine precipitates are precipitated particularly in the surface layer portion in combination with Ti at the time of the post-heat treatment, thereby improving the fatigue resistance in the plastic strain region. When the content of C is less than 0.19%, it is difficult to obtain the desired strength (YS. gtoreq.770 MPa) and fatigue resistance in the plastic strain region. On the other hand, if the content of C exceeds 0.50%, the toughness and weldability of the steel member, for example, steel pipe, cannot be ensured, so the upper limit is set. It is further preferable that the content of C is more than 0.28%. Further, the content of C is preferably 0.30% or less.
Si：0.002～1.5％
Si is an element that secures a desired strength by solid-solution strengthening and improves fatigue resistance in a plastic strain region. When the content of Si is less than 0.002%, the strength is insufficient. On the other hand, if the content exceeds 1.5%, weldability deteriorates. Therefore, the content of Si is preferably limited to 0.002 to 1.5%. It is more preferable that the content of Si is 0.05% or more. Further, the content of Si is preferably 0.35% or less.
Mn：0.4～2.5％
Mn has an effect of securing a desired strength by low-temperature transformation strengthening at the time of post-heat treatment, and improving fatigue resistance in a plastic strain region. When the Mn content is less than 0.4%, the effect is not sufficiently exhibited, while when the Mn content exceeds 2.5%, weldability deteriorates. Therefore, the Mn content is preferably limited to 0.4 to 2.5%. It is more preferable that the Mn content is 1.09% or more. Further, the Mn content is preferably 1.99% or less.
Al：0.01～0.19％
Al is a deoxidizing element in steel making, and has the effects of suppressing the growth of austenite grains in a hot rolling process, making the grains fine, obtaining a desired ferrite grain size (1 to 50 μm) after post-heat treatment, and improving fatigue resistance in a plastic strain region. If the Al content is less than 0.01%, these effects cannot be obtained, and the ferrite grain size becomes coarse, while if the Al content exceeds 0.19%, weldability decreases, and fatigue resistance tends to decrease due to the increase in oxide inclusions. It is more preferable that the Al content is 0.041% or more. Further, the Al content is preferably 0.080% or less.
Cr：0.001～0.90％
Cr has an effect of securing a desired strength by low-temperature transformation strengthening at the time of post-heat treatment, and improving fatigue resistance in a plastic strain region. When the content of Cr is less than 0.001%, the effect is not sufficiently exhibited, while when the content of Cr exceeds 0.90%, weldability deteriorates. Therefore, the content of Cr is preferably limited to 0.001 to 0.90%. The content of Cr is more preferably 0.001 to 0.19%.
B：0.0001～0.0050％
B has an effect of securing a desired strength by low-temperature transformation strengthening at the time of post-heat treatment, and improving fatigue resistance in a plastic strain region. If the content of B is less than 0.0001%, the effect is not sufficiently exhibited, while if the content of B exceeds 0.0050%, the fatigue resistance tends to be lowered. Therefore, the content of B is preferably limited to 0.0001 to 0.0050%. Further, the content of B is preferably 0.0005% or more. Further, the content of B is preferably 0.0035% or less.
P: less than 0.019% (including 0%)
P reduces fatigue resistance in the plastic strain region and deteriorates resistance weldability by co-segregation with solidification of Mn. When the content of P exceeds 0.019%, the adverse effect becomes remarkable, and therefore, 0.019% is preferably set as the upper limit.
S: less than 0.015% (including 0%)
S forms MnS and the like and exists as inclusions in steel, and deteriorates fatigue resistance as a starting point of fatigue crack in a plastic strain region. When the content of S exceeds 0.015%, the adverse effect becomes significant, and therefore, it is preferable to set the upper limit to 0.015%. It is further preferable that the S content is 0.005% or less.
N: less than 0.008% (including 0%)
N forms TiN with Ti, precipitates as coarse precipitates, and consumes solid-solution Ti. Thus, N reduces the following effects: by adding Ti, 0.005% or more of Ti is present as solid-solution Ti at the stage of the raw hot-rolled steel sheet, and by heat treatment after forming, 0.005% or more of Ti is precipitated as fine precipitates having a grain size of 20nm or less, and thereby fatigue resistance in a plastic strain region which is particularly excellent can be obtained. When the content of N exceeds 0.008%, the adverse effect becomes significant, and therefore, it is preferable to set 0.008% as the upper limit. It is more preferable that the content of N is 0.0049% or less.
O: less than 0.003% (including 0%)
O exists as oxide inclusions, and deteriorates the fatigue resistance of steel. When the content of O exceeds 0.003%, the adverse effect becomes remarkable, and therefore, it is preferable to set the upper limit to 0.003%. The O content is more preferably 0.002% or less.
Sn: less than 0.10% (including 0%)
Sn exists as a solid solution element, and decreases the hot ductility of steel. When the content of Sn exceeds 0.10%, the adverse effect becomes significant, and therefore, it is preferable to set 0.10% as the upper limit. It is further preferable that the Sn content is 0.03% or less.
The balance being Fe and unavoidable impurities. In the present invention, the following elements may be added for the purpose of further improving the effects of the present invention.
Nb：0.001～0.15％
Nb is precipitated as carbide, and has the effect of suppressing grain growth of recovery and recrystallization in the hot rolling step to obtain a desired ferrite grain size (1 to 50 μm), and may be contained as necessary. When the content of Nb is less than 0.001%, these effects cannot be obtained. On the other hand, if the content of Nb exceeds 0.15%, coarse precipitates precipitate in the surface layer portion due to strain-induced precipitation during hot rolling, fine precipitates in the surface layer portion decrease, and fatigue resistance in the plastic strain region decreases, so 0.15% is set as the upper limit. Therefore, when Nb is contained, the Nb content is set to 0.001 to 0.15%. The content of Nb is more preferably 0.001 to 0.009%.
V：0.001～0.15％
V is precipitated as carbide, has an effect of suppressing grain growth of recovery and recrystallization in the hot rolling step to obtain a desired ferrite grain size (1 to 50 μm), and may be contained as necessary. When the content of V is less than 0.001%, these effects cannot be obtained. On the other hand, if the content of V exceeds 0.15%, coarse precipitates precipitate in the surface layer portion due to strain-induced precipitation during hot rolling, fine precipitates in the surface layer portion decrease, and fatigue resistance in the plastic strain region decreases, so 0.15% is set as the upper limit. Therefore, when V is contained, the content of V is set to 0.001 to 0.15%. Further, the content of V is preferably 0.001 to 0.049%.
W：0.001～0.15％
W precipitates as carbide, has an effect of supplementing the effect of suppressing the grain growth of recovery and recrystallization in the hot rolling step to obtain a desired ferrite grain size (1 to 50 μm), and may be contained as necessary. When the content of W is less than 0.001%, these effects cannot be obtained. On the other hand, when the W content exceeds 0.15%, coarse precipitates precipitate in the surface layer portion due to strain-induced precipitation during hot rolling, fine precipitates in the surface layer portion decrease, and fatigue resistance in the plastic strain region decreases, so 0.15% is set as the upper limit. Therefore, when W is contained, the content of W is set to 0.001 to 0.15%. The content of W is more preferably 0.001 to 0.049%.
Mo：0.001～0.45％
Mo has an effect of supplementing the effect of securing a desired strength by low-temperature transformation strengthening or precipitation strengthening at the time of post-heat treatment and improving fatigue resistance in a plastic strain region, and may be contained as necessary. When the content of Mo is less than 0.001%, the effect is not exhibited, while when the content of Mo exceeds 0.45%, weldability deteriorates. Therefore, when Mo is contained, the content of Mo is set to 0.001 to 0.45%. The content of Mo is more preferably 0.001 to 0.30%.
Cu：0.001～0.45％、Ni：0.001～0.45％
Cu and Ni are elements having an effect of supplementing the effect of Mn to improve fatigue strength and an effect of improving corrosion resistance of the steel material, and may be contained in each of Cu and Ni as necessary. These effects are exhibited when Cu and Ni are contained in amounts of 0.001% or more, respectively, and when Cu and Ni are contained in amounts exceeding 0.45%, respectively, the weldability is lowered, so that 0.45% is set as the upper limit, respectively. Therefore, when Cu is contained, the content of Cu is set to 0.001 to 0.45%. In addition, when Ni is contained, the content of Ni is set to 0.001-0.45%. It is more preferable that all of the elements are 0.35% or less.
Ca：0.0001～0.005％
Ca has a so-called morphology controlling effect of Ca (al) s (o) for forming elongated MnS into grains, and has an effect of suppressing the occurrence of fatigue cracks and improving fatigue resistance, and may be contained as necessary. This effect is exhibited when the content is 0.0001% or more, and when the content exceeds 0.005%, the fatigue resistance is rather lowered due to the increase in the size of the nonmetallic inclusions, so that 0.005% is set as the upper limit. Therefore, when Ca is contained, the content of Ca is set to 0.0001 to 0.005%.
Sb：0.0001～0.10％
Sb is preferentially segregated on the surface, has an action of suppressing the intrusion of N from the atmosphere in the hot rolling step or the post-heat treatment step, and suppressing the decrease in the effect of adding B due to the formation of BN, and may be contained as necessary. Although this effect is exhibited when the content is 0.0001% or more, the effect is saturated even when the content exceeds 0.10%, and therefore 0.10% is set as the upper limit. Therefore, when Sb is contained, the content of Sb is set to 0.0001 to 0.10%. The content of Sb is more preferably 0.0001 to 0.030%.
Further, the steel member of the present invention is preferably: the post-heat-treated steel sheet has a structure in which the ferrite phase having a grain size of 200 [ mu ] m from the surface to the thickness direction has an average crystal grain size of 1 to 50 [ mu ] m and Ti carbide having a grain size of 1.0 to 20nm is precipitated in the ferrite phase having a grain size of 200 [ mu ] m from the surface to the thickness direction, and the difference (absolute value) between the average hardness of 200 [ mu ] m from the surface to the thickness direction and the average hardness in the vicinity of the thickness center excluding the center segregation portion is Δ HV50 point or less in terms of Vickers Hardness (HV).
The microstructure of the steel member, the precipitated state of the precipitates, and the sectional hardness are important in ensuring excellent fatigue resistance in the plastic strain region. When the average crystal grain size of the ferrite phase of 200 μm in the thickness direction from the surface after the post heat treatment exceeds 50 μm, the initial fatigue crack occurs rapidly and largely, and it is difficult to ensure the fatigue resistance in the desired plastic strain region. On the other hand, it is industrially and economically difficult to make the average crystal grain size of the ferrite phase after the post-heat treatment smaller than 1 μm, and therefore the lower limit is set.
The ferrite phase referred to herein is a matrix iron of a body-centered cubic lattice, and includes polygonal ferrite, acicular ferrite, widmannstatten ferrite, bainitic ferrite, bainite, and a low-carbon (C content of 1% or less) martensite structure. The second phase other than the ferrite phase includes austenite, carbide, pearlite, and high-carbon martensite (the C content exceeds 1%).
The structure of the steel member of the present invention preferably has the ferrite phase as a main phase. The main phase is a phase occupying not less than 51% by volume, preferably not less than 80% by volume, and may be 100%.
In addition, in order to ensure surface hardness and to ensure fatigue resistance in a high plastic strain region, the Ti carbide size in the ferrite phase of 200 μm from the surface to the thickness direction is important. By precipitating Ti carbide of 1.0 to 20nm in a ferrite phase of 200 μm in the thickness direction from the surface, the occurrence of fatigue-initiating cracks can be suppressed, and the size thereof can be reduced, whereby the fatigue resistance in an excellent plastic strain region can be further improved. The amount of Ti carbide precipitates of 1.0 to 20nm is not particularly limited. In addition, Ti carbides having different sizes are allowed to precipitate in addition to Ti carbides having a size of 1.0 to 20 nm.
It is important to ensure excellent fatigue resistance in the plastic strain region that the difference between the average hardness of 200 μm from the surface in the thickness direction and the average hardness in the vicinity of the thickness center excluding the center segregation portion is Δ HV50 or less. When the difference between the average hardness of 200 μm from the surface in the thickness direction and the average hardness in the vicinity of the thickness center excluding the center segregation portion exceeds Δ HV50 point, initial fatigue cracks occur early and largely, and it is difficult to ensure fatigue resistance in a desired plastic strain region. Therefore, the difference between the average hardness of 200 μm from the surface in the thickness direction and the average hardness in the vicinity of the thickness center excluding the center segregation portion is preferably Δ HV50 point or less.
The difference between the average hardness of 200 μm from the surface to the thickness direction and the average hardness in the vicinity of the thickness center except the center segregation portion was measured as follows: the micro Vickers hardness (HV (0.1)) was measured at a pitch of 25 μm in the thickness direction between 50 and 200 μm in the thickness direction by a load of 0.1kgf, and the average of 7 points was determined to obtain the value HV (0.1)SHV (0.1) at 7 points was measured at a pitch of 25 μm in the thickness direction with the center of the thickness of the plate kept as the center and avoiding the center segregation portion, and averaged to obtain a value HV (0.1)CDetermination of HV (0.1)SAnd HV (0.1)CThe difference is HV (0.1)C-HV(0.1)S。
(Hot rolled Steel sheet of stock)
The hot-rolled steel sheet for a steel member of the present invention (raw hot-rolled steel sheet) is particularly suitable for obtaining the steel member of the present invention.
The raw material hot rolled steel sheet of the present invention contains 0.031-0.200% by mass of Ti, and 0.005% or more of Ti is present in the structure as solid solution Ti. Thus, after the forming process and after the predetermined heat treatment, 0.005% or more of Ti can be precipitated in the structure of the steel member as fine precipitates having a grain size of 20nm or less, and the steel member excellent in fatigue resistance and further excellent in strength characteristics in the plastic strain region can be obtained.
The raw hot-rolled steel sheet of the present invention has the same composition as the steel member.
In the hot-rolled steel sheet as a raw material according to the present invention, it is preferable that the thicknesses of both the front end portion and the rear end portion in the longitudinal direction are 5 to 50% thicker than the thickness of the intermediate portion (the longitudinal direction central portion) other than both the longitudinal direction end portions. Thus, the effect of improving fatigue resistance in the plastic strain region of the welded portion in the case of cutting the hot-rolled steel sheet as a material into a predetermined width and then welding the cut pieces in the longitudinal direction for use is enhanced, as in the case of a coiled tubing.
(production method)
Next, a method for manufacturing a steel member and a hot-rolled steel sheet as a material thereof according to the present invention will be described. In the following description, the temperature is set to the surface temperature of a billet or the like unless otherwise specified.
In the present invention, a billet obtained by casting a steel having the above-described composition is used as a starting material. The method for producing the starting material is not particularly limited, and examples thereof include a method in which the molten steel having the above composition is melted by a usual melting method such as a converter and is formed into a billet by a usual casting method such as a continuous casting method.
First, a method for producing a hot-rolled steel sheet (raw hot-rolled steel sheet) as a material of a steel member according to the present invention will be described.
The raw material hot-rolled steel sheet of the present invention is produced by hot-rolling a steel slab containing 0.031 to 0.200% of Ti under predetermined conditions.
Higher than log ([ Ti-N × 48 ÷ 14)][C])＝-7000/(TTiCalculated equilibrium solid solubility temperature T of (DEG C.) +273) +2.75TiUnder the temperature condition of
The billet extraction temperature in the hot rolling step affects the size of precipitates and the amount of solid-solution Ti after hot rolling by the re-solid solution and precipitation state of Ti in the steel, and is important for ensuring good fatigue resistance after post-heat treatment. The extraction temperature is the equilibrium solid solubility temperature T calculated by the following formula (1)TiIn the following case, coarse Ti precipitated during continuous casting remains as undissolved carbonitride, and the amount of solid-solution Ti is less than 0.005% at the stage of the raw hot-rolled steel sheet, and the fatigue resistance in the plastic strain region which is particularly excellent after post-heat treatment cannot be obtained. By heating at a temperature higher than the equilibrium solid solubility temperature T calculated from (1) belowTiExtracting a slab under the temperature condition of (1) to obtain a hot rolled steel sheetIn the stage (2), 0.005% or more of Ti is present as solid-solution Ti, and 0.005% or more of Ti can be precipitated as fine precipitates having a grain diameter of 20nm or less by heat treatment after forming, and fatigue resistance in a plastic strain region which is particularly excellent can be obtained. It is further preferable that the billet extraction temperature is 1620K or less from the viewpoint of preventing coarsening of the crystal grain size, and the soaking time of the billet (at a temperature higher than the equilibrium solid solution temperature T) is more preferable from the viewpoint of ensuring the uniformity of the solid solution state of Ti and a sufficient solid solution timeTiThe temperature of) is preferably 10 minutes or more.
log([Ti-N×48÷14][C])＝-7000/(TTi(℃)+273)+2.75…(1)
Wherein Ti, N and C in the formula (1) represent the content (mass%) of each element in the steel slab.
TTi-finish rolling temperature of 400 ℃ or higher
The temperature of the hot rolling and finish rolling is lower than TTiAt-400 ℃, the amount of solid-solution Ti present at the stage of the raw hot-rolled steel sheet, particularly in the vicinity of the surface (within 200 μm from the surface and back surface), is less than 0.005% due to additional shear strain by the upper and lower rolls in the vicinity of the surface or strain-induced precipitation induced by heat dissipation by the rolls and cooling water, and no particularly excellent fatigue resistance in the plastic strain region can be obtained after post-heat treatment. By setting the finish rolling temperature TTiAt least 400 ℃ and at least 0.005% of Ti is present as solid solution Ti including the vicinity of the surface at the stage of the hot rolled raw steel sheet, and 0.005% or more of Ti can be precipitated as fine precipitates having a grain size of 20nm or less by heat treatment after forming, and fatigue resistance in a plastic strain region which is particularly excellent can be obtained.
At the slave TTi-400 ℃ to TTiCooling at an average cooling rate of 10 ℃/sec or more in a temperature range of-500 DEG C
From TTi-400 ℃ to TTiWhen the average cooling rate in the temperature range of-500 ℃ is less than 10 ℃/s, TiC is precipitated from the hot rolling output to the coiling and is hot rolled in the raw materialThe amount of solid-solution Ti present at the stage of the steel sheet is less than 0.005%, and fatigue resistance in a plastic strain region which is particularly excellent after post-heat treatment cannot be obtained. By following TTi-400 ℃ to TTiRapidly cooling at an average cooling rate of 10 ℃/sec or more in a temperature range of-500 ℃, wherein at the stage of the raw hot-rolled steel sheet, 0.005% or more of Ti is present as solid-solution Ti including the vicinity of the surface, and 0.005% or more of Ti can be precipitated as fine precipitates having a grain size of 20nm or less by heat treatment after forming, whereby fatigue resistance characteristics in a plastic strain region which are particularly excellent can be obtained.
TTiCoiling temperature below-500 DEG C
Coiling temperature exceeding TTiAt-500 ℃, precipitation of Ti precipitates is promoted until the coil is cooled, the amount of solid-solution Ti existing at the stage of the raw hot-rolled steel sheet is less than 0.005%, and fatigue resistance in a plastic strain region which is particularly excellent after post-heat treatment cannot be obtained. By setting the coiling temperature to TTiAt the stage of hot-rolled steel sheet of-500 ℃ or lower, 0.005% or more of Ti is present as solid-solution Ti including the vicinity of the surface, and 0.005% or more of Ti can be precipitated as fine precipitates having a grain size of 20nm or less by heat treatment after forming, whereby fatigue resistance in a plastic strain region which is particularly excellent 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 determined from the surface temperatures.
By the above-described production method, a hot-rolled steel sheet (raw hot-rolled steel sheet) in which 0.005% or more of Ti is present in the structure as solid-solution Ti can be obtained.
Next, a method for manufacturing a steel member of the present invention will be described.
The steel member of the present invention is produced by subjecting the hot-rolled steel sheet as the material to forming and then to a predetermined heat treatment. The forming process is not particularly limited, and for example, when the steel member is a steel pipe, a pipe-making process is exemplified. When the steel member is a welded steel pipe, the welding process may be performed after the pipe-making process. For example, when the steel member is a formed member such as an automobile structural member, press working or the like is exemplified.
After the forming process, heat treatment was performed under the following conditions.
Heating to a temperature higher than 550 ℃ and lower than 1050 ℃, and cooling at an average cooling rate of 10 ℃/sec or higher within a temperature range of 550-400 DEG C
By performing the heat treatment after the forming process of the raw hot-rolled steel sheet, heating the raw hot-rolled steel sheet to a temperature exceeding 550 ℃ and 1050 ℃ or lower, and then cooling the sheet at an average cooling rate of 10 ℃/sec or higher in a temperature range of 550 to 400 ℃, 0.005% or more of Ti is precipitated as fine precipitates having a particle size of 20nm or less, and fatigue resistance in a plastic strain region which is particularly excellent can be obtained. When the heating temperature is 550 ℃ or lower, the solid-solution Ti does not precipitate as fine precipitates of 20nm or less, and the fatigue resistance in the plastic strain region which is particularly excellent cannot be obtained. When the heating temperature exceeds 1050 ℃, the particle size of the ferrite phase exceeds 50 μm, and it is difficult to obtain particularly excellent fatigue resistance in the plastic strain region. In addition, when the cooling rate is less than 10 ℃/sec in the temperature range of 550 to 400 ℃, sufficient strength (YS ≥ 770MPa) cannot be obtained. The heating temperature is more preferably in the range of 700 to 1000 ℃.
Further, although not particularly limited, for example, in the case of manufacturing a welded steel pipe, a raw hot-rolled steel sheet is processed by one or more of pickling, cold rolling, annealing, and plating while maintaining a black skin, if necessary, then slit into a predetermined sheet width, coils are welded and joined to each other in the longitudinal direction by 1 coil or more, the coils are formed into a substantially circular cross-section by roll forming or press forming, the ends are joined by a method such as high-frequency resistance welding or laser welding, heated to a temperature exceeding 550 ℃ and 1050 ℃ or less in an on-line or off-line manner, and then cooled at an average cooling rate of 10 ℃/sec or more within a temperature range of 550 to 400 ℃ to form a steel pipe wound into a coil.
For example, in the case of manufacturing a formed member, the raw hot-rolled steel sheet is punched out to a predetermined size while being kept in a black state or after being subjected to any one or more of pickling, cold rolling, annealing, and plating as necessary, and is then heated to a temperature exceeding 550 ℃ and 1050 ℃ or lower after being formed into a member, and then cooled at an average cooling rate of 10 ℃/sec or more within a temperature range of 550 to 400 ℃. Thus, 0.005% or more of Ti is precipitated as fine precipitates having a particle diameter of 20nm or less, and fatigue resistance in a plastic strain region is remarkably excellent.
Examples
(example 1)
Slabs having compositions (steel types C to L) shown in Table 1 were extracted from a heating furnace at slab surface temperatures of about 1220 ℃ and slab center temperatures of about 1210 ℃ so as to obtain a finish rolling reduction of 91%, a finish rolling temperature at the center of the coil width of about 860 ℃, a minimum finish rolling temperature in the coil width direction of about 850 ℃, and a pass temperature T of the steel sheetTi-400 ℃ to TTiAnd hot rolling at an average cooling rate of about 20 ℃/sec at a temperature of-500 ℃ to a coiling temperature of about 560 ℃ to obtain a hot rolled steel sheet (thickness: about 5mm, thickness of front and rear end portions about 10% relative to the thickness of the longitudinal central portion) (Nos. 3 to 12).
Hot-rolled raw steel sheets (nos. 2 and 13) were produced in the same manner as described above except that slabs having the compositions (steel type a) shown in table 1 were extracted from the heating furnace at slab surface temperatures of about 1250 ℃ and slab center temperatures of about 1245 ℃, and slabs having the compositions (steel type B, M) shown in table 1 were extracted from the heating furnace at slab surface temperatures of about 1335 ℃ and slab center temperatures of about 1335 ℃ to produce hot-rolled raw steel sheets at coil width center portion finish rolling temperatures of about 940 ℃.
Next, these hot rolled steel sheets as raw materials were pickled, slit-processed into predetermined width dimensions, continuously formed into open pipes, and resistance-welded by high-frequency resistance welding to obtain welded steel pipes having a width reduction ratio of 4%, an outer diameter of 50.8mm, and a wall thickness of about 5 mm. The welded steel pipe as a whole was continuously subjected to high-frequency heating, and the following heat treatment was performed: heating at 920 ℃ for about 5 seconds, spray-cooling the outer surface with water, and cooling at an average cooling rate of about 50 ℃/sec within a temperature range of 550 to 400 ℃.
Test pieces were cut out from these welded steel pipes, and a structure observation test, a quantitative test of precipitates and solid solution amount, a tensile test, a plastic strain region fatigue test, and a low-temperature toughness test were performed. The test method is as follows.
(1) Tissue observation test
A structure observation test piece was cut out so that the circumferential direction cross section of these welded steel pipes was an observation surface, the structure was observed by a scanning Electron microscope (3000 times) after polishing and nital etching, and the average grain size of the ferrite phase was determined by an EBSD (Electron back scattering Diffraction) method with the angle of inclination with the adjacent crystal grains of 15 ° or more as the grain boundary. As the average particle diameter from the surface to 200 μm in the thickness direction, a value obtained by measuring 3 points at a pitch of 50 μm between 50 and 200 μm in the thickness direction and averaging the measured values, and a value obtained by measuring 3 points at a pitch of 50 μm in the thickness direction while avoiding the center segregation portion with the center of the thickness center as the center and averaging the measured values, were obtained, respectively.
(2) Quantitative test of amount of precipitate and solid solution
A specimen piece of 20mm X30 mm size was cut out of the welded steel pipe, and the cut piece was immersed in 10% AA electrolyte (10 vol% acetylacetone-1 mass% tetramethylammonium chloride-methanol) at 20mA/cm2Constant current electrolysis was carried out for about 0.2 g. The sample piece after electrolysis with the deposit adhering to the surface was taken out from the electrolytic solution, immersed in a sodium hexametaphosphate aqueous solution (500mg/l) (hereinafter referred to as a SHMP aqueous solution), and subjected to ultrasonic vibration to peel the deposit from the sample piece and extract the deposit into the SHMP aqueous solution. Then, the precipitate-containing SHMP aqueous solution was filtered through a filter in the order of pore diameters of 100nm and 20nm, the residue on the filter and the filtrate after filtration were analyzed by an ICP emission spectrometer, and the absolute amounts of Ti in the residue on the filter and in the filtrate were measured to obtain precipitates having a particle diameter of more than 100nm and precipitates having a particle diameter of 100nm or less and more than 20nmThe precipitates of (4) and the absolute amounts of Ti, Tilp, and Tisp contained in the precipitates having a particle diameter of 20nm or less. The electrolytic mass was determined by measuring the mass of the sample piece after the precipitates were peeled off and subtracting the mass from the mass of the sample piece before the electrolysis.
With respect to Ti in a solid solution state (solid solution Ti), the concentration of Ti and Fe as a comparative element in the solution was measured by an ICP mass spectrometry using the electrolytic solution after electrolysis as an analysis solution. Based on the obtained concentrations, the concentration ratios of Ti to Fe are calculated, and the contents of Ti in a solid solution state are obtained by multiplying the concentrations by the contents of Fe in the sample. The content of Fe in the sample can be determined by subtracting the total of the composition values other than Fe from 100%. The quantitative test of the amount of precipitates and solid solutions was conducted not only on the welded steel pipe after the post heat treatment but also on the welded steel pipe before the post heat treatment.
(3) Tensile test
Test pieces of JIS No.12 were cut out from these welded steel pipes in accordance with the specifications of JIS Z2201 with the L direction as the stretching direction, and a tensile test was performed in accordance with the specifications of JIS Z2241 to obtain tensile characteristics (tensile strength TS, yield strength YS, total elongation El).
(4) Fatigue test in plastic strain region
After flattening, a plate-like L-direction fatigue test piece having a cross-sectional dimension of a parallel portion of about 5mm in plate thickness × 5mm in plate width and 12mm in parallel portion length was cut out from the welded steel pipe, and a fatigue test was performed under the conditions of a tensile mode, a strain control mode, a strain ratio of 0, a total strain range of 2.0%, and a cycle number of 0.125 Hz. The number of cycles at which the maximum tensile load was reduced by 25% from the initial load was determined as the number of repetitions until fracture.
(5) Low temperature toughness test
The steel pipes were developed so that the pipe length direction (L direction) was the length of the test piece, charpy test pieces (2mmV notch, 1/2 size) were cut out of these welded steel pipes in accordance with the specification of JIS Z2202, and charpy impact test was performed in accordance with the specification of JIS Z2242 to determine the fracture transition temperature and evaluate the low-temperature toughness.
Further, the average hardness (HV (0.1)) of 200 μm in the thickness direction from the surface was measured by the above-mentioned methodS) Average hardness in the vicinity of the center of the plate thickness except for the center segregation portion (HV (0.1)C) The difference DeltaHV (HV (0.1) between the average hardness of 200 μm from the surface in the thickness direction and the average hardness near the thickness center except the center segregation portion was determinedC-HV(0.1)S)。
The obtained results are shown in table 2.
Figure BDA0002224904880000221
Figure BDA0002224904880000231
In the present invention examples (Nos. 1 to 11), the number of cycles in the plastic strain region fatigue test was 1000 cycles or more, and the fatigue resistance in the plastic strain region was excellent. In the present invention examples, YS was 770MPa or more, and all of them were excellent in strength characteristics. In the examples of the present invention, the Charpy fracture transition temperature was-30 ℃ or lower, and the low-temperature toughness was excellent. On the other hand, No.12 in which the steel composition does not satisfy the range of the present invention and Ti precipitated as precipitates having a grain size of 20nm or less is less than 0.005%, and No.13 in which the steel composition does not satisfy the range of the present invention do not obtain the fatigue resistance in the desired plastic strain region.
(example 2)
A steel slab having a composition of steel grade A, B, C shown in Table 1 was hot-rolled under the conditions shown in Table 3 to obtain a hot-rolled steel sheet (thickness: about 5mm, thickness of front and rear end portions about 10% thicker than a longitudinal central portion). Next, these hot rolled steel sheets as raw materials were pickled, slit-processed into a predetermined width, continuously formed into an open pipe, and resistance-welded to the open pipe by high-frequency resistance welding to obtain a welded steel pipe having a width reduction ratio of 4%, an outer diameter of 50.8mm, and a wall thickness of about 5 mm. The welded steel pipe as a whole was continuously subjected to high-frequency heating and heat treatment under the conditions shown in table 3. Test pieces were cut out from these welded steel pipes, and subjected to a structure observation test, a quantitative test for precipitates and solid solution amount, a tensile test, a plastic strain region fatigue test, a low-temperature toughness test, and a vickers hardness measurement.
In the case of No.23, a hot-rolled raw steel sheet was pickled, punched out to a predetermined size, press-worked to form a formed member, and heat-treated under the conditions shown in table 3. Then, a test piece was cut out from the molded part, and each of the above tests was performed.
The obtained results are shown in table 4. The results of the above Nos. 1 to 3 are also shown in tables 3 and 4.
Figure BDA0002224904880000251
Figure BDA0002224904880000261
In the present invention examples (Nos. 21 to 23), the number of cycles in the plastic strain region fatigue test was 1000 cycles or more, and the fatigue resistance in the plastic strain region was excellent. In the present invention examples, YS was 770MPa or more, and all of them were excellent in strength characteristics. In the examples of the present invention, the Charpy fracture transition temperature was-30 ℃ or lower, and the low-temperature toughness was excellent. On the other hand, in Nos. 14 to 20 in which the Ti content precipitated as precipitates having a particle diameter of 20nm or less is outside the range of the present invention, the fatigue resistance in the desired plastic strain region was not obtained.

Claims (8)

1. A steel member having a composition containing, in mass%, C: 0.19 to 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 to 0.0050%, Ti: 0.031-0.200%, P: less than 0.019% (including 0%), S: 0.015% or less (including 0%), N: less than 0.008% (including 0%), O: 0.003% or less (including 0%), Sn: 0.10% or less (including 0%), and the balance of Fe and unavoidable impurities, wherein 0.005% or more of Ti is precipitated in the structure as precipitates having a particle diameter of 20nm or less.
2. The steel member as claimed in claim 1, further comprising, in mass%, on the basis of the composition, an additive selected from the group consisting of Nb: 0.001-0.15%, V: 0.001-0.15%, W: 0.001 to 0.15%, Mo: 0.001 to 0.45%, Cu: 0.001-0.45%, Ni: 0.001-0.45%, Ca: 0.0001 to 0.005%, Sb: 0.0001-0.10% of one or more than two.
3. A steel member as claimed in claim 1 or 2, wherein the steel member is a welded steel pipe.
4. A hot-rolled steel sheet for a steel member according to any one of claims 1 to 3, wherein the composition is as defined in claim 1 or 2, and 0.005% or more of Ti is present in the structure as solid-solution Ti.
5. A hot-rolled steel sheet for a steel member according to claim 4, wherein the thickness of each of the front end portion and the rear end portion, which are both longitudinal end portions, is 5 to 50% thicker than the thickness of the central portion in the longitudinal direction.
6. A method for manufacturing a steel member according to any one of claims 1 to 3, wherein a hot-rolled steel sheet having the composition defined in claim 1 or 2 and having a structure in which 0.005% or more of Ti is present as solid-solution Ti is subjected to forming, and then subjected to heat treatment as follows: heating to a temperature higher than 550 ℃ and lower than 1050 ℃, and then cooling at an average cooling rate of 10 ℃/sec or higher within a temperature range of 550-400 ℃.
7. A method of manufacturing a steel member as claimed in claim 6 wherein there will beA steel slab having said composition defined in claim 1 or 2 and having a temperature higher than the equilibrium solid solubility temperature T calculated by the following formula (1)TiAfter billet extraction at the temperature of (1), at TTiFinishing the finish rolling at a temperature of-400 ℃ or higher, at a temperature of from TTi-400 ℃ to TTiCooling at a temperature in the range of-500 ℃ at an average cooling rate of 10 ℃/s or more, at TTiCoiling at a temperature of-500 ℃ or lower to thereby manufacture the hot rolled steel sheet,
log([Ti-N×48÷14][C])＝-7000/(TTi(℃)+273)+2.75…(1)
wherein Ti, N and C in the formula (1) represent the content (mass%) of each element in the steel slab.
8. A method for producing a hot-rolled steel sheet for a steel member, which is the method for producing a hot-rolled steel sheet for a steel member according to claim 4 or 5, wherein a steel slab having the composition defined in claim 1 or 2 is set to a temperature higher than the equilibrium solid solubility temperature T calculated by the following formula (1)TiAfter billet extraction at the temperature of (1), at TTiFinishing the finish rolling at a temperature of-400 ℃ or higher, at a temperature of from TTi-400 ℃ to TTiCooling at a temperature in the range of-500 ℃ at an average cooling rate of 10 ℃/s or more, at TTiCoiling at a temperature below-500 ℃,
log([Ti-N×48÷14][C])＝-7000/(TTi(℃)+273)+2.75…(1)
wherein Ti, N and C in the formula (1) represent the content (mass%) of each element in the steel slab.
CN201880023595.9A 2017-04-07 2018-03-29 Steel member, hot-rolled steel sheet for said steel member, and method for producing same Active CN110494582B (en)

Applications Claiming Priority (3)

Application Number Priority Date Filing Date Title
JP2017-076411 2017-04-07
JP2017076411 2017-04-07
PCT/JP2018/013076 WO2018186273A1 (en) 2017-04-07 2018-03-29 Steel member, hot-rolled steel sheet for said steel member and production methods therefor

Publications (2)

Publication Number Publication Date
CN110494582A CN110494582A (en) 2019-11-22
CN110494582B true CN110494582B (en) 2021-08-03

Family

ID=63713115

Family Applications (1)

Application Number Title Priority Date Filing Date
CN201880023595.9A Active CN110494582B (en) 2017-04-07 2018-03-29 Steel member, hot-rolled steel sheet for said steel member, and method for producing same

Country Status (7)

Country Link
US (1) US20200190618A1 (en)
JP (1) JP6631715B2 (en)
KR (1) KR102319579B1 (en)
CN (1) CN110494582B (en)
CA (1) CA3057814C (en)
MX (1) MX2019011941A (en)
WO (1) WO2018186273A1 (en)

Families Citing this family (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN110494583B (en) * 2017-04-07 2021-10-26 杰富意钢铁株式会社 Steel member, hot-rolled steel sheet for said steel member, and method for producing same
US20220090224A1 (en) * 2019-01-09 2022-03-24 Nippon Steel Corporation Hot-rolled steel sheet and weld joint, and methods for producing same
CN113265581B (en) * 2020-02-17 2022-10-21 宝山钢铁股份有限公司 Steel pipe for stabilizer bar and manufacturing method thereof
KR20220147727A (en) * 2020-09-10 2022-11-03 닛폰세이테츠 가부시키가이샤 Steel plate and its manufacturing method

Family Cites Families (21)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP3491339B2 (en) 1994-06-17 2004-01-26 住友金属工業株式会社 Manufacturing method of high tension electric resistance welded steel pipe
DE10323693B3 (en) * 2003-05-22 2004-09-09 Muhr Und Bender Kg Sheet element from flexible rolled strip for vehicle components is formed into variable wall thickness over its length, producing tubes or strip of non-round cross section
US20080247900A1 (en) * 2004-07-16 2008-10-09 Jfe Steel Corporation Component for Machine Structure, Method of Producing the Same and Material for Induction Hardening
JP4819489B2 (en) * 2005-11-25 2011-11-24 Ｊｆｅスチール株式会社 High strength steel plate with excellent uniform elongation characteristics and method for producing the same
JP4466619B2 (en) * 2006-07-05 2010-05-26 Ｊｆｅスチール株式会社 High tensile welded steel pipe for automobile structural members and method for manufacturing the same
KR20150127739A (en) * 2007-05-06 2015-11-17 누코 코포레이션 A thin cast strip product with microalloy additions, and method for making the same
JP5196934B2 (en) 2007-09-27 2013-05-15 日新製鋼株式会社 High fatigue life quenched and tempered steel pipe and method for manufacturing the same
JP5359296B2 (en) * 2008-01-17 2013-12-04 Ｊｆｅスチール株式会社 High strength steel plate and manufacturing method thereof
WO2009118945A1 (en) * 2008-03-26 2009-10-01 新日本製鐵株式会社 Hot rolled steel sheet possessing excellent fatigue properties and stretch-flange ability and process for producing the hot rolled steel sheet
CN102341521B (en) * 2009-05-27 2013-08-28 新日铁住金株式会社 High-strength steel sheet, hot-dipped steel sheet, and alloy hot-dipped steel sheet that have excellent fatigue, elongation, and collision characteristics, and manufacturing method for said steel sheets
CN101713046B (en) * 2009-12-14 2013-09-18 钢铁研究总院 Preparation method of superfine grain martensitic steel reinforced and controlled by nano precipitated phase
CN102191438A (en) * 2010-03-18 2011-09-21 宝山钢铁股份有限公司 Steel plate for seamless high pressure gas cylinder, and manufacture method thereof
US20120132322A1 (en) * 2010-11-30 2012-05-31 Kennametal Inc. Abrasion resistant steel, method of manufacturing an abrasion resistant steel and articles made therefrom
US20150004050A1 (en) 2012-01-18 2015-01-01 Jfe Steel Corporation Steel strip for coiled tubing and method of manufacturing the same
CN102676927A (en) * 2012-06-12 2012-09-19 钢铁研究总院 High-Ti micro-alloyed medium thickness steel plate and preparation method thereof
CN103320684A (en) * 2013-07-18 2013-09-25 王艳林 Steel multielement second-phase thermal equilibrium solid solution numerical analysis model
KR101716624B1 (en) * 2013-09-10 2017-03-14 가부시키가이샤 고베 세이코쇼 Method for manufacturing press-molded article, and press-molded article
JP6528522B2 (en) * 2015-04-17 2019-06-12 日本製鉄株式会社 High strength hot rolled steel sheet excellent in ductility, fatigue characteristics and corrosion resistance and manufacturing method thereof
JP6252692B2 (en) * 2015-07-27 2017-12-27 Ｊｆｅスチール株式会社 High strength hot rolled steel sheet and method for producing the same
JP6668662B2 (en) * 2015-09-30 2020-03-18 日本製鉄株式会社 Steel sheet excellent in fatigue characteristics and formability and method for producing the same
JP6699307B2 (en) * 2016-04-08 2020-05-27 日本製鉄株式会社 Hot-rolled steel sheet and its manufacturing method

Also Published As

Publication number Publication date
JPWO2018186273A1 (en) 2019-04-11
MX2019011941A (en) 2019-11-28
KR102319579B1 (en) 2021-10-29
US20200190618A1 (en) 2020-06-18
WO2018186273A1 (en) 2018-10-11
CA3057814A1 (en) 2018-10-11
CA3057814C (en) 2022-07-12
CN110494582A (en) 2019-11-22
JP6631715B2 (en) 2020-01-15
KR20190125397A (en) 2019-11-06

Similar Documents

Publication Publication Date Title
CN110832098B (en) Hot-rolled steel sheet and method for producing same
JP6048580B2 (en) Hot rolled steel sheet and manufacturing method thereof
EP2816132B1 (en) Steel sheet, plated steel sheet, method for producing steel sheet, and method for producing plated steel sheet
JP4288201B2 (en) Manufacturing method of automotive member having excellent hydrogen embrittlement resistance
CN110494582B (en) Steel member, hot-rolled steel sheet for said steel member, and method for producing same
KR100802237B1 (en) Spring steel with Excellent Resistance to Hydrogen Embrittlement and Steel wire and Spring obtained from the steel
JP5251208B2 (en) High-strength steel sheet and its manufacturing method
EP1693476A1 (en) Steel product for structural member of automobile and method for production thereof
WO2009125874A1 (en) High-strength steel sheets which are extremely excellent in the balance between burring workability and ductility and excellent in fatigue endurance, zinc-coated steel sheets, and processes for production of both
TWI468530B (en) Cold rolled steel plate, plated steel plate, and method of manufacturing the same
CN114438418A (en) Hot-formed member and method for manufacturing same
KR20190037680A (en) Steel wire rod and steel wire for spring having corrosion fatigue resistance and method of manufacturing thereof
KR102544884B1 (en) High-strength hot-dip galvanized steel sheet and manufacturing method thereof
JP6172399B2 (en) High strength steel plate and manufacturing method thereof
CN113383097A (en) Steel sheet, method for producing steel sheet, and plated steel sheet
JP2006009116A (en) Steel sheet for hot pressing
JP2008163409A (en) High tensile strength welded steel pipe for automobile structural member, and method for producing the same
JP2007291464A (en) High-strength steel material and its production method
CN110494583B (en) Steel member, hot-rolled steel sheet for said steel member, and method for producing same
KR102658163B1 (en) High-strength steel plate and its manufacturing method
JP5499560B2 (en) High tensile steel material for automobile undercarriage members having excellent formability and torsional fatigue resistance and method for producing the same
JP2009203492A (en) High-tensile welded steel pipe for automobile structural member, and method for producing the same
JP2001181794A (en) High strength spring steel
CN113544300A (en) High-strength steel sheet and method for producing same
JPH1177116A (en) Manufacture of high-fatigue-strength steel tube

Legal Events

Date Code Title Description
PB01 Publication
PB01 Publication
SE01 Entry into force of request for substantive examination
SE01 Entry into force of request for substantive examination
GR01 Patent grant
GR01 Patent grant