WO2010150450A1 - 高強度極細鋼線及びその製造方法 - Google Patents

高強度極細鋼線及びその製造方法 Download PDF

Info

Publication number
WO2010150450A1
WO2010150450A1 PCT/JP2010/002859 JP2010002859W WO2010150450A1 WO 2010150450 A1 WO2010150450 A1 WO 2010150450A1 JP 2010002859 W JP2010002859 W JP 2010002859W WO 2010150450 A1 WO2010150450 A1 WO 2010150450A1
Authority
WO
WIPO (PCT)
Prior art keywords
steel wire
mass
ferrite phase
concentration
wire
Prior art date
Application number
PCT/JP2010/002859
Other languages
English (en)
French (fr)
Japanese (ja)
Inventor
高橋淳
小坂誠
児玉順一
樽井敏三
鈴木環輝
Original Assignee
新日本製鐵株式会社
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by 新日本製鐵株式会社 filed Critical 新日本製鐵株式会社
Priority to CN2010800018607A priority Critical patent/CN102066599B/zh
Priority to JP2010540974A priority patent/JP4719320B2/ja
Priority to US12/994,169 priority patent/US20110168302A1/en
Priority to EP10778818.4A priority patent/EP2447382B1/en
Publication of WO2010150450A1 publication Critical patent/WO2010150450A1/ja

Links

Images

Classifications

    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/52—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length
    • C21D9/525—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for wires; for strips ; for rods of unlimited length for wire, for rods
    • 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
    • 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
    • C21D8/06—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of rods or wires
    • C21D8/065—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of rods or wires of ferrous 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
    • C22C38/02—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
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • 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
    • 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
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel with boron
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/009—Pearlite
    • 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
    • C21D7/00—Modifying the physical properties of iron or steel by deformation
    • C21D7/02—Modifying the physical properties of iron or steel by deformation by cold working

Definitions

  • the present invention relates to a high-strength steel wire used for a steel cord or saw wire of an automobile tire and a manufacturing method thereof. More specifically, the present invention relates to an ultra fine steel wire having a wire diameter of 0.04 to 0.4 mm and a strength of 4500 MPa class or higher, which is cold-drawn and strengthened by using a die.
  • This application claims priority based on Japanese Patent Application No. 2009-148051 filed in Japan on June 22, 2009, the contents of which are incorporated herein by reference.
  • High strength of steel wire is accompanied by a decrease in ductility, and it is important to suppress this. It is also important that high ductility steel wire is not substantially deteriorated by aging because its properties deteriorate due to aging at room temperature (20 ° C to 40 ° C, several days to several years). is there.
  • High-strength steel wires are generally manufactured by drawing a wire having a pearlite structure using a die or the like. This processing reduces the pearlite lamella spacing, and introduces a large amount of dislocations in the ferrite phase, thereby increasing the tensile strength. It has recently been clarified that cementite in a pearlite structure is refined and decomposed when the wire drawing strain becomes very large. However, since the structure is particularly fine, the relationship between the location and state of these carbons and the mechanical properties has not been clarified, and there are many unclear points regarding the cause of ductile deterioration. In an actual high-strength steel wire, the structure and local strain in the steel wire do not always appear to be the same in the surface region and the central region, which also affects the properties of the steel wire. it is conceivable that.
  • Patent Document 1 Patent Document 2, and Patent Document 3 define chemical components such as C, Si, Mn, and Cr, respectively, as conventional knowledge of a high-strength means with little reduction in ductility.
  • a high-strength, high-ductility, high-carbon steel wire rod for ultrafine wires has been proposed.
  • the maximum tensile strength of the steel wire is 3500 to 3600 MPa, and there is a limit to increasing the strength of the ultrafine steel wire.
  • Patent Document 4 proposes a high-strength steel high-toughness steel wire rod in which the chemical composition, non-metallic inclusion structure, and area fraction of proeutectoid cementite are controlled. Furthermore, Patent Document 5 discloses a method for producing a high-strength steel and a high toughness ultrafine wire steel that controls the chemical composition of the steel and the area reduction rate of the final die. However, even with these techniques, it has been impossible to realize an ultrafine steel wire having a tensile strength of 4500 MPa or more and high ductility.
  • Patent Document 9 specifies the variation in carbon concentration.
  • Patent Document 10 defines the degree of difference in lamella spacing that affects the variation in carbon concentration. However, these describe the overall variation and do not prescribe the carbon concentration at a specific location.
  • patent document 11 shows the manufacturing method of the steel wire for obtaining a favorable characteristic by prescribing
  • Patent Document 12 and Patent Document 13 define a range of residual stress from the viewpoint of fatigue resistance and longitudinal crack resistance.
  • the residual compressive stress is preferable, the absolute value of the value is small, and the range for obtaining a very good balance between ductility and strength has not yet been defined.
  • the ductility of the high-strength ultrafine steel wire is responsible for the ductility of the ferrite phase. If the ductility of the ferrite phase is maintained, the ductility is ensured even at high strength. However, when the wire drawing strain increases, cementite is generally decomposed, C atoms diffuse into the ferrite phase, and the carbon concentration in the ferrite phase increases. In Non-Patent Document 1, in a cold-rolled steel sheet, when the carbon concentration in the ferrite phase increases, dynamic strain aging occurs in which the dislocations in the ferrite phase are fixed by carbon during the tensile test, resulting in a significant decrease in ductility. It is stated to cause.
  • the present invention provides a high-strength steel wire, particularly a high-strength ultrafine steel wire, having a high strength of 4500 MPa or more and excellent ductility, against the background of the current situation as described above.
  • the present invention employs the following means in order to solve the above problems.
  • the chemistry of C 0.7 to 1.2% by mass, Si: 0.05 to 2.0% by mass, Mn: 0.2 to 2.0% by mass
  • a steel wire containing a component and the balance containing Fe and unavoidable impurities, the steel wire has a pearlite structure, and the average C concentration of the ferrite phase central portion of the outermost layer of the steel wire is 0.2.
  • the steel wire has a residual compressive stress of 600 MPa or more in the longitudinal direction of the steel wire of the outermost layer.
  • the steel wire according to the above (1) or (2) may be a high-strength ultrafine steel wire having a tensile strength of 4500 MPa or more.
  • the high-strength ultrafine steel wire described in the above (3) may be a steel cord.
  • the high-strength ultrafine steel wire described in (3) above may be a saw wire.
  • a second aspect of the present invention is a method for producing a steel wire having a tensile strength of 4500 MPa or more, wherein C: 0.7 to 1.2% by mass, Si: 0.05 to 2.0
  • a patenting step of producing a pearlite structure by performing a patenting treatment on a steel wire containing a chemical component of mass%, Mn: 0.2 to 2.0 mass%, and the balance containing Fe and inevitable impurities;
  • a residual stress applying step for applying compressive stress is a method for producing a steel wire having a tensile strength of 4500 MPa or more, wherein C: 0.7 to 1.2% by mass, Si: 0.05 to 2.0
  • a patenting step of producing a pearlite structure by performing a patenting treatment on a steel wire
  • the steel wire according to the present invention can exhibit high strength and ductility because the carbon concentration of the ferrite phase central portion of the outermost layer of the steel wire having a pearlite structure is controlled and the residual compressive stress is applied. it can. Moreover, since it becomes possible to provide the high strength steel wire which has sufficient ductility and tensile strength, the weight reduction of a product is attained.
  • the present inventors have found that the carbon (hereinafter referred to as C) concentration in the ferrite phase in the outermost layer of the steel wire in a strongly processed drawn pearlite structure It has been newly found that the residual stress in the longitudinal direction of the steel wire outermost layer strongly affects the ductility of the steel wire. This is thought to be because, in bending and twisting, the outermost layer of the steel wire is subjected to a stronger stress than the inside and becomes the starting point of fracture.
  • C carbon
  • the outermost layer of the steel wire is subjected to severer processing on the inside of the steel wire, and undergoes severe temperature changes due to frictional heat generation and the like. Therefore, the structure and state are clearly different from the inside of the steel wire. Accordingly, cementite decomposition further proceeds, and the C concentration in the ferrite phase of the outermost layer is generally higher than the C concentration in the ferrite phase inside the steel wire. Since the outermost layer of the steel wire is the most strongly involved in the properties, it has been found that a steel wire excellent in balance between strength and ductility can be realized almost by controlling the structure of the outermost layer.
  • a high-strength steel wire is generally obtained by strengthening a wire having a pearlite structure by applying a high-drawing process using a die or the like.
  • a phenomenon in which cementite in the pearlite structure is refined and decomposed and C is dissolved in the ferrite phase due to high wire drawing strain generated during high wire drawing is generally obtained.
  • the present inventors have developed a three-dimensional atom probe method (hereinafter referred to as 3DAP) that can measure the local concentration of C in a miniaturized region and a technique for producing a needle sample from the outermost steel wire layer that has been made possible for the first time.
  • 3DAP three-dimensional atom probe method
  • the relationship between the C concentration in the ferrite phase and the strength and ductility of the steel wire at every location in the steel wire was examined in detail.
  • the ductility may be significantly reduced, particularly when the C concentration in the ferrite phase of the steel wire surface layer is increased, or when the residual stress of the same outermost layer is pulled or weakly compressed in the longitudinal direction of the steel wire. It was located (see FIG. 1).
  • the average C concentration in the ferrite phase central portion of the outermost steel wire layer is set to a specific value or less, and the residual in the longitudinal direction of the steel wire on the surface It was concluded that the stress should be a sufficiently large compressive stress.
  • the present inventors prepared samples having a tensile strength of 4500 MPa or more by various production methods, the tensile strength and ductility, the average C concentration of the ferrite phase center of the surface pearlite structure, and the surface residual stress. I investigated the relationship with. The average C concentration at the ferrite phase outermost layer of the steel wire was measured by 3DAP, and the residual stress was examined by X-ray diffraction. Tensile strength measurement was performed with a tensile tester, and a torsion test, which is one of ductility evaluations, was performed with a torsion tester. As a ductility index, the number of twists to break was measured.
  • FIG. 1 shows the average C concentration of the ferrite phase center at a position of 1 ⁇ m below the surface of the steel wire, the residual stress in the steel wire longitudinal direction of the outermost layer of the steel wire, and the ductility expressed by the number of twists until breaking by the torsion test.
  • the result of having investigated about the relationship of is shown.
  • samples with 20 or more twists were indicated by white circles (good ductility), and samples with 25 or more twists were indicated by white squares (very good ductility). Samples of less than 20 times are indicated by black triangles (poor ductility).
  • a steel wire having a tensile strength of 4500 MPa or more and good ductility is a large compression with an average C concentration of 0.2 mass% or less in the ferrite phase center portion of the outermost layer of the steel wire and a residual stress of -600 MPa or less. Only observed when. Furthermore, a steel wire with very good ductility was observed when the average C concentration in the ferrite phase center was 0.1% by mass or less and the residual stress was a strong compressive stress of ⁇ 600 MPa or less. .
  • the average C concentration in the ferrite phase outermost layer of the steel wire is 0.2 mass% or less, more preferably 0.1 mass or less,
  • the residual stress in the longitudinal direction of the steel wire outermost layer is desirably ⁇ 600 MPa, more preferably ⁇ 700 MPa or less.
  • the average C concentration is preferably as low as possible, but the carbon concentration in the center of the ferrite phase of the pearlite structure of the final patenting material is in principle the lowest carbon concentration. Therefore, the lower limit value of the average C concentration at the center of the ferrite phase of the outermost layer may be set to 0.0001% by mass.
  • the maximum value of the residual compressive stress corresponds in principle to the yield stress of the steel wire, but may be substantially ⁇ 3000 MPa. Applying a compressive stress larger than this leads to a significant cost increase and is not practical.
  • the steel wire outermost layer refers to a region within a depth of 2 ⁇ m from the surface, excluding the plating phase and the surface heterogeneous phase.
  • the ferrite phase central portion of the pearlite structure of the outermost steel wire layer is a region including a distance of 1/4 of the width of the ferrite phase on both sides from the position of the central surface of the ferrite phase (region of half the width of the ferrite phase). ).
  • a steel wire according to an embodiment of the present invention based on the above discovery has a C content of 0.7 to 1.2 mass%, a Si content of 0.05 to 2.0 mass%, and a Mn content of 0.2 to 2.0 mass%. It is a steel wire that is contained by mass and the balance contains Fe and inevitable impurities.
  • This steel wire has a drawn pearlite structure, the average C concentration of the ferrite phase central portion of the outermost layer is 0.2% by mass or less, and the residual in the longitudinal direction of the steel wire of the outermost layer of the steel wire
  • the compressive stress is 600 MPa or more. The reason for limitation will be described in detail below. “%” Shown below means “% by mass” unless otherwise specified.
  • C has an effect of increasing the tensile strength after the patenting treatment and increasing the wire drawing work hardening rate, and it becomes possible to increase the tensile strength with less wire drawing distortion. If the C content is 0.7% or less, it will be difficult to achieve the intended high strength steel wire in the present invention. On the other hand, if it exceeds 1.2%, the pro-eutectoid cementite will enter the austenite grain boundaries during the patenting process. It precipitates and the wire drawing workability deteriorates, causing breakage during wire drawing. For this reason, the range of C content is defined as 0.7 to 1.2%.
  • Si is an effective element for strengthening the ferrite phase in pearlite and for deoxidation of steel. If the Si content is less than 0.05%, the above effect cannot be expected. On the other hand, if it exceeds 2%, hard SiO 2 inclusions harmful to the wire drawing workability are likely to be generated. For this reason, the range of Si content is prescribed to 0.05 to 2.0%.
  • Mn is not only necessary for deoxidation and desulfurization, but is also an effective element for improving the hardenability of steel and increasing the tensile strength after patenting.
  • Mn content is less than 0.2%, the above effect cannot be obtained.
  • it exceeds 2.0% the above effect is saturated, and further, the pearlite transformation during the patenting process is completed. The processing time becomes too long and productivity is lowered. For this reason, the range of the Mn content is defined as 0.2 to 2.0%.
  • the above-described steel wire according to an embodiment of the present invention may further include one or more of Cr, Ni, V, Nb, Mo, and B for the following reasons.
  • Cr increases the tensile strength after the patenting treatment by reducing the interval between the pearlite cementite phases and also improves the wire drawing work hardening rate.
  • the Cr content is less than 0.05%, the effect of the above action is small.
  • the Cr content exceeds 1.0%, the pearlite transformation finish time during the patenting process becomes long and the productivity is lowered. For this reason, it is preferable to keep the Cr content in the range of 0.05 to 1.0%.
  • Ni has the effect of making pearlite produced by transformation during the patenting process to have good wire drawing workability. However, if the Ni content is less than 0.05%, the above effect cannot be obtained, and 1.0% Even if the amount exceeds 50%, the effect corresponding to the amount added is small. For this reason, it is preferable to keep the Ni content in the range of 0.05 to 1.0%.
  • V has the effect of increasing the tensile strength during patenting by reducing the interval between the pearlite cementite phases, but this effect is insufficient when the V content is less than 0.01%. If it exceeds 5%, the effect is saturated. For this reason, it is preferable to keep the V content in the range of 0.01 to 0.5%.
  • Nb like V, has the effect of reducing the cementite phase interval and increasing the tensile strength during the patenting treatment, but if the Nb content is less than 0.001%, it is insufficient. If it exceeds%, the effect is saturated. For this reason, it is preferable to keep the Nb content in the range of 0.001 to 0.1%.
  • Mo Mo, like V, has the effect of increasing the cementite phase interval and increasing the tensile strength during the patenting treatment. However, if the Mo content is less than 0.01%, it is insufficient. If it exceeds%, the effect is saturated. For this reason, it is preferable to keep the Mo content in the range of 0.01 to 0.1%.
  • B has an effect of fixing N as BN and preventing aging deterioration due to N, and in order to fully exhibit this effect, the B content in the steel material needs to be contained by 0.0001% or more. is there. On the other hand, even if it is added so that the B content in the steel material exceeds 0.01%, the effect is saturated, and B content beyond this is not preferable because it causes the production cost to be increased. For this reason, in the present invention, when B is contained in the steel material, the B content is preferably within the range of 0.0001 to 0.01%.
  • impurities inevitably mixed in the manufacturing process may be contained, but it is preferable that impurities are not mixed as much as possible.
  • Skin pass drawing which is one of important manufacturing methods, is a method of drawing with a particularly small area reduction than the normal area reduction (10% or more).
  • the area reduction rate is preferably 1% to 6%, and more preferably 2% to 5%. If the area reduction is less than 1%, it becomes difficult to process the entire surface of the steel wire, and if it exceeds 7%, the amount of processing is too large, and preferable compressive residual stress and ferrite phase on the surface. The carbon concentration in can not be obtained.
  • This skin pass wire drawing may be performed independently by a single die method, or may be performed simultaneously with normal wire drawing by a double die method.
  • Shot peening is performed after wire drawing. Shot peening is a method of irradiating the entire steel wire with a spherical shot of a specific size for a specific time at a specific pressure, and creating a processed layer or a strained layer only in the surface region of the steel wire. Shot peening is, for example, an air projection method, an air pressure of 4 to 5 ⁇ 10 5 Pa, a time of 5 to 10 seconds is preferable, and a shot spherical shape of 10 to 100 ⁇ m is preferable. It is effective to irradiate a sufficient amount of the entire surface of the steel wire.
  • the final stage drawing speed is 200 m / min or less, preferably 50 m / min or less.
  • a die having an approach angle of 8 to 12 ° and a dynamic friction coefficient of 0.1, preferably 0.05 or less is used in the final stage including the skin pass and the previous drawing step.
  • (Group C manufacturing method) C1 After the wire drawing, heating and holding at 60 to 300 ° C. is performed at 0.1 to 24 hours, more preferably at 180 to 260 ° C. for 20 seconds to 15 minutes. Due to aging during or after wire drawing, cementite decomposes and supersaturated carbon dissolved in the ferrite phase is discharged, reducing the carbon concentration in the ferrite phase. However, when this temperature is too high, spherical cementite or transition carbide is formed, and when it is too low, the effect is small. It is necessary to set a temperature suitable for the steel material type and wire drawing conditions.
  • a process with a large area reduction rate of 20% or more is inserted once, preferably a plurality of times during wire drawing excluding the three stages before the final stage.
  • the C concentration in the ferrite phase in the steel wire can be accurately measured by the three-dimensional atom probe method (3DAP).
  • 3DAP three-dimensional atom probe method
  • FIB focused ion beam
  • solute C concentration may show different values depending on the position in the ferrite phase.
  • the C concentration at the ferrite phase / cementite phase interface position is high, and the value is the smallest at the ferrite phase center position.
  • the average C concentration of a region (a region that is a half of the width of the ferrite phase) including a distance of 1/4 of the width of the ferrite phase on both sides from the position of the center plane of the ferrite phase is defined.
  • the C concentration in the ferrite phase including the ferrite phase / cementite phase interface By analyzing with 3DAP, it is possible to measure the C concentration in the ferrite phase including the ferrite phase / cementite phase interface, so by selecting and cutting out a box of a specific size in the region to be examined from the measurement data, By calculating the ratio of C atoms to all atoms in the box, the C concentration in the ferrite phase can be obtained in atomic%. This can be converted to mass% by multiplying by 12/56. Such measurement was performed for a plurality of ferrite phase center portions to obtain an average, and this was taken as the average C concentration of the ferrite phase center portion.
  • FIGS. 2A to 2F show a method of preparing a needle sample for measuring the C concentration of the ferrite phase center within 1 ⁇ m from the surface of the steel wire
  • FIG. 3 shows the use of the prepared needle sample. The C distribution measured by 3DAP and the C concentration at the center of the ferrite phase are shown.
  • FIG. 2A In order to produce a needle sample in a region 1 ⁇ m from the surface of the steel wire, for example, as shown in FIG. 2A, a rod-like block including the steel wire surface on one side is cut out from the steel wire surface region by FIB.
  • This block is fixed on the needle base as shown in FIG. 2B by using vapor deposition (deposition) of tungsten or the like, for example.
  • this block As shown in FIG. 2C, this block is processed by FIB so that the tip end portion is thin.
  • FIG. 2D is a diagram of the processed block observed from above, and it can be seen that the tip has a rod shape including the steel wire surface. Then, the tip part was processed into a needle shape by irradiating a ring-shaped beam from the upper part.
  • FIG. 2D is a diagram of the processed block observed from above, and it can be seen that the tip has a rod shape including the steel wire surface. Then, the tip part was processed into a needle shape by
  • FIG. 2F is a view of the needle sample thus made, observed from the side.
  • the needle tip position was prepared so as to correspond to the inside of 1 ⁇ m from the steel wire surface.
  • a needle sample of the outermost steel wire layer can be prepared.
  • the dark portion indicates that the C concentration is high and the light portion indicates that the C concentration is low. Accordingly, a dark band-like region indicates a cementite phase that has undergone wire drawing, and a lightly colored region between them indicates a ferrite phase that has undergone wire drawing. It is shown that C is also dissolved in the ferrite phase.
  • the carbon concentration at the center of the ferrite phase can be estimated.
  • the C concentration is 0.18% by mass.
  • the ferrite phase center is located in the middle of the two cementite phases, and includes a region that includes up to a quarter of the width of the ferrite phase on both sides from the center surface of the ferrite phase (region that is half the width of the ferrite phase). ).
  • the width of the ferrite phase is not necessarily constant depending on the processing amount and the location of the sample, and there is a region of 10 nm or less in a narrow portion. If the cementite region is included in the box position, it becomes higher than the actual C concentration in the ferrite phase. Therefore, the box position to be analyzed was the ferrite phase center, and the box width was half the width of the ferrite phase. Further, the average C concentration is estimated by averaging the measured values of the C concentration at the central portion of 5 or more, preferably 10 or more different ferrite phases.
  • the residual stress in the outermost steel wire layer can be measured with high accuracy by, for example, the X-ray diffraction method.
  • it can be measured accurately by the Debye fitting method using a micro-region X-ray diffractometer capable of measuring a local region.
  • This method is a method of fitting the reflection of the crystal grain of the steel wire as a Debye ring and examining the magnitude direction of the residual stress from the Debye distortion.
  • the depth region including the surface is determined from the penetration depth of X-rays. For example, when the X-ray source is Cr, an integrated value with a depth of several ⁇ m on the surface can be obtained.
  • Another method for examining the residual stress on the surface of the steel wire is a melting method (Hane method) as needed.
  • test material having the chemical composition shown in Table 1 After making the test material having the chemical composition shown in Table 1 into a predetermined wire diameter by hot rolling, performing a patenting treatment and a wire drawing using a lead bath, so that the tensile strength becomes 4500 MPa or more, A high-strength ultrafine wire steel made of a drawn pearlite structure with a brass plating having a wire diameter of 0.04 to 0.40 mm was produced. Brass plating was performed after pickling after the final patenting treatment.
  • Table 2 shows the true strain, the manufacturing method, the wire diameter, the average C concentration of the ferrite phase in the outermost layer of the steel wire, the residual stress in the outermost layer of the steel wire, the tensile strength, and the torsion test. Shows the number of twists to break.
  • the manufacturing method is represented by symbols indicating the contents described above.
  • the torsion test the test piece was fixed at a gripping distance 100 times the diameter of both ends of the test piece, and the number of torsion until breaking was examined. When the tensile strength was 4500 MPa or more and the number of twists was 20 times or more, the ductility was evaluated as good, and when the tensile strength was 25 times or more, the ductility was evaluated as very good.
  • the C concentration in the ferrite phase of the steel wire outermost layer was measured at the surface of 1 ⁇ m by 3DAP using the method described above, and the residual stress in the longitudinal direction of the steel wire outermost layer was measured by the Debyling fitting method described above. .
  • the residual stress is negative, it represents compressive stress, and when it is positive, it represents tensile stress.
  • test no. 1 to 6 are examples of the present invention, and others are comparative examples. As can be seen from the table, all of the inventive examples have a tensile strength of 4500 MPa or more, an average C concentration in the center of the ferrite phase of the outermost layer of 0.2% by mass or less, and a residual stress of ⁇ 600 MPa or less ( The residual compressive stress is 600 MPa or more. As a result, an ultrafine steel wire having a sufficient ductility with a high number of twists can be realized. In particular, test no. In Nos. 1 and 2, the number of twists was 25 or more, which was very good.
  • test no. 7 to 20 are comparative examples, and the tensile strength was 4500 MPa or more, but the number of twists was insufficient.
  • No. 7 to 9 are comparative examples in which the components of the steel wire are outside the scope of the present invention.
  • No. 7 since the amount of C in the steel wire was too small, and the amount of wire drawing strain was increased, the C concentration at the ferrite phase central portion exceeded the specified value, and the ductility decreased.
  • No. 8 is the amount of Si in the steel wire.
  • 9 is a comparative example in which the amount of C is higher than the range of the present invention. In these comparative examples, the residual stress and the C concentration at the ferrite phase center were within the specified range, but the ductility was lowered.
  • No. Nos. 10 to 13 are comparative examples in which the composition and residual stress of the steel wire are within the scope of the present invention, but the C concentration at the ferrite phase center of the outermost layer is not less than a specified value. In these comparative examples, the ductility decreased.
  • No. Nos. 14 to 16 are comparative examples in which the component of the steel wire and the C concentration in the ferrite phase central portion are within the range of the present invention, but the residual stress is out of the range. In these comparative examples, the ductility decreased.
  • No. Nos. 17 to 20 are comparative examples in which the C concentration and the residual stress at the ferrite phase center of the outermost layer are both out of the range. In these comparative examples, the ductility decreased.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Materials Engineering (AREA)
  • Mechanical Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Manufacturing & Machinery (AREA)
  • Heat Treatment Of Steel (AREA)
  • Ropes Or Cables (AREA)
  • Heat Treatment Of Strip Materials And Filament Materials (AREA)
PCT/JP2010/002859 2009-06-22 2010-04-21 高強度極細鋼線及びその製造方法 WO2010150450A1 (ja)

Priority Applications (4)

Application Number Priority Date Filing Date Title
CN2010800018607A CN102066599B (zh) 2009-06-22 2010-04-21 高强度极细钢线及其制造方法
JP2010540974A JP4719320B2 (ja) 2009-06-22 2010-04-21 高強度極細鋼線及びその製造方法
US12/994,169 US20110168302A1 (en) 2009-06-22 2010-04-21 High-strength ultrathin steel wire and method of manufacturing the same
EP10778818.4A EP2447382B1 (en) 2009-06-22 2010-04-21 High-strength ultra-fine steel wire and manufacturing method therefor

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
JP2009148051 2009-06-22
JP2009-148051 2009-06-22

Publications (1)

Publication Number Publication Date
WO2010150450A1 true WO2010150450A1 (ja) 2010-12-29

Family

ID=43386240

Family Applications (1)

Application Number Title Priority Date Filing Date
PCT/JP2010/002859 WO2010150450A1 (ja) 2009-06-22 2010-04-21 高強度極細鋼線及びその製造方法

Country Status (7)

Country Link
US (1) US20110168302A1 (zh)
EP (1) EP2447382B1 (zh)
JP (1) JP4719320B2 (zh)
KR (1) KR101260598B1 (zh)
CN (1) CN102066599B (zh)
TW (1) TWI412608B (zh)
WO (1) WO2010150450A1 (zh)

Cited By (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN102965546A (zh) * 2012-11-20 2013-03-13 无锡康柏斯机械科技有限公司 一种帘线机
WO2016024635A1 (ja) * 2014-08-15 2016-02-18 新日鐵住金株式会社 伸線加工用鋼線
JP2017141494A (ja) * 2016-02-10 2017-08-17 新日鐵住金株式会社 強度と延性のバランスに優れた高強度極細鋼線
KR101830537B1 (ko) 2016-11-07 2018-02-20 주식회사 포스코 피로 저항성이 우수한 고강도 강선 및 이의 제조방법
WO2018203504A1 (ja) * 2017-05-02 2018-11-08 横浜ゴム株式会社 ビードリング及びその製造方法
CN112223569A (zh) * 2020-09-28 2021-01-15 王佩 一种耐磨线切割复合线材及其制备方法
WO2024024401A1 (ja) * 2022-07-29 2024-02-01 住友電気工業株式会社 鋼線、及び鋼線の製造方法

Families Citing this family (14)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2008069409A (ja) * 2006-09-14 2008-03-27 Bridgestone Corp 高強度高炭素鋼線およびその製造方法
KR101316198B1 (ko) * 2011-11-15 2013-10-08 주식회사 포스코 고연성 선재, 강재 및 그 제조방법
EP2993246B1 (en) * 2013-04-30 2019-03-20 Nippon Steel & Sumitomo Metal Corporation Flat steel wire
WO2015119241A1 (ja) 2014-02-06 2015-08-13 新日鐵住金株式会社 フィラメント
CN105960477B (zh) * 2014-02-06 2018-11-06 新日铁住金株式会社 钢线
KR101870063B1 (ko) * 2014-04-24 2018-06-22 신닛테츠스미킨 카부시키카이샤 고강도 스틸 코드용 필라멘트
WO2015163407A1 (ja) 2014-04-24 2015-10-29 新日鐵住金株式会社 高強度スチールコード用線材
CN105063508B (zh) * 2015-08-26 2017-01-18 武汉钢铁(集团)公司 含硼高端帘线钢及其生产方法
JP6481770B2 (ja) * 2015-10-23 2019-03-13 新日鐵住金株式会社 伸線加工用鋼線材
KR102303599B1 (ko) * 2017-05-18 2021-09-23 닛폰세이테츠 가부시키가이샤 선재, 강선, 및 강선의 제조 방법
CN109811565A (zh) * 2017-11-21 2019-05-28 江苏法尔胜技术开发中心有限公司 压实喷丸面股耐磨钢丝绳及其生产方法
CN110238235A (zh) * 2019-06-11 2019-09-17 安徽环新集团有限公司 用于汽车发动机的卡簧线材生产线及生产方法
CN113699438B (zh) * 2021-07-20 2022-07-08 武汉钢铁有限公司 一种86级低成本帘线钢及其制备工艺
CN113913681B (zh) * 2021-08-26 2022-03-25 武汉钢铁有限公司 一种高强度低断丝率的帘线钢、轧制方法及其用途

Citations (13)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS60204865A (ja) 1984-03-28 1985-10-16 Kobe Steel Ltd 高強度で高靭延性の極細線用高炭素鋼線材
JPS6324046A (ja) 1986-07-16 1988-02-01 Kobe Steel Ltd 高靭性高延性極細線用線材
JPH0323674A (ja) 1989-06-21 1991-01-31 Mitsubishi Electric Corp 半導体装置
JPH06145895A (ja) 1992-10-30 1994-05-27 Kobe Steel Ltd 高強度高靭性鋼線材、該鋼線材を用いた極細鋼線およびその製法並びに撚り鋼線
JPH07113119A (ja) 1993-10-15 1995-05-02 Kobe Steel Ltd 高強度高靭延性極細鋼線の製法
JPH11199979A (ja) 1998-01-16 1999-07-27 Nippon Steel Corp 疲労特性の優れた高強度極細鋼線およびその製造方法
JPH11199980A (ja) 1998-01-20 1999-07-27 Nippon Steel Corp 高強度極細鋼線
JP2001220649A (ja) 2000-02-08 2001-08-14 Nippon Steel Corp 延性及び疲労特性の優れた高強度極細鋼線
JP2001279381A (ja) 2000-03-29 2001-10-10 Kobe Steel Ltd 耐縦割れ性に優れた高強度高靭性極細鋼線及びその製造方法
JP2003334606A (ja) 2003-03-31 2003-11-25 Nippon Steel Corp 高強度極細鋼線の製造方法
JP2006249561A (ja) 2005-03-14 2006-09-21 Nippon Steel Corp 延性に優れた高強度極細鋼線
JP2007262496A (ja) 2006-03-28 2007-10-11 Nippon Steel Corp 延性に優れた高強度極細鋼線
JP2008208450A (ja) 2007-01-30 2008-09-11 Nippon Steel Corp 強度延性バランスに優れた高強度極細鋼線の製造方法

Family Cites Families (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPH08232046A (ja) * 1995-02-23 1996-09-10 Nippon Steel Corp 耐捻回割れ性に優れた高強度鋼線
US6790294B1 (en) * 1999-02-19 2004-09-14 Suncall Corporation Spring with excellent fatigue endurance property and surface treatment method for producing the spring
JP2001279380A (ja) * 2000-03-29 2001-10-10 Sumitomo Electric Ind Ltd 高靭性鋼線およびその製造方法
US6949149B2 (en) * 2002-12-18 2005-09-27 The Goodyear Tire & Rubber Company High strength, high carbon steel wire

Patent Citations (13)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS60204865A (ja) 1984-03-28 1985-10-16 Kobe Steel Ltd 高強度で高靭延性の極細線用高炭素鋼線材
JPS6324046A (ja) 1986-07-16 1988-02-01 Kobe Steel Ltd 高靭性高延性極細線用線材
JPH0323674A (ja) 1989-06-21 1991-01-31 Mitsubishi Electric Corp 半導体装置
JPH06145895A (ja) 1992-10-30 1994-05-27 Kobe Steel Ltd 高強度高靭性鋼線材、該鋼線材を用いた極細鋼線およびその製法並びに撚り鋼線
JPH07113119A (ja) 1993-10-15 1995-05-02 Kobe Steel Ltd 高強度高靭延性極細鋼線の製法
JPH11199979A (ja) 1998-01-16 1999-07-27 Nippon Steel Corp 疲労特性の優れた高強度極細鋼線およびその製造方法
JPH11199980A (ja) 1998-01-20 1999-07-27 Nippon Steel Corp 高強度極細鋼線
JP2001220649A (ja) 2000-02-08 2001-08-14 Nippon Steel Corp 延性及び疲労特性の優れた高強度極細鋼線
JP2001279381A (ja) 2000-03-29 2001-10-10 Kobe Steel Ltd 耐縦割れ性に優れた高強度高靭性極細鋼線及びその製造方法
JP2003334606A (ja) 2003-03-31 2003-11-25 Nippon Steel Corp 高強度極細鋼線の製造方法
JP2006249561A (ja) 2005-03-14 2006-09-21 Nippon Steel Corp 延性に優れた高強度極細鋼線
JP2007262496A (ja) 2006-03-28 2007-10-11 Nippon Steel Corp 延性に優れた高強度極細鋼線
JP2008208450A (ja) 2007-01-30 2008-09-11 Nippon Steel Corp 強度延性バランスに優れた高強度極細鋼線の製造方法

Non-Patent Citations (1)

* Cited by examiner, † Cited by third party
Title
JOURNAL OF JAPAN INSTITUTE OF METAL, vol. 45, no. 9, 1981, pages 942 - 947

Cited By (11)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
CN102965546A (zh) * 2012-11-20 2013-03-13 无锡康柏斯机械科技有限公司 一种帘线机
WO2016024635A1 (ja) * 2014-08-15 2016-02-18 新日鐵住金株式会社 伸線加工用鋼線
JPWO2016024635A1 (ja) * 2014-08-15 2017-06-22 新日鐵住金株式会社 伸線加工用鋼線
US10329646B2 (en) 2014-08-15 2019-06-25 Nippon Steel & Sumitomo Metal Corporation Steel wire for drawing
JP2017141494A (ja) * 2016-02-10 2017-08-17 新日鐵住金株式会社 強度と延性のバランスに優れた高強度極細鋼線
KR101830537B1 (ko) 2016-11-07 2018-02-20 주식회사 포스코 피로 저항성이 우수한 고강도 강선 및 이의 제조방법
WO2018203504A1 (ja) * 2017-05-02 2018-11-08 横浜ゴム株式会社 ビードリング及びその製造方法
JP2018188764A (ja) * 2017-05-02 2018-11-29 横浜ゴム株式会社 ビードリング及びその製造方法
CN112223569A (zh) * 2020-09-28 2021-01-15 王佩 一种耐磨线切割复合线材及其制备方法
WO2024024401A1 (ja) * 2022-07-29 2024-02-01 住友電気工業株式会社 鋼線、及び鋼線の製造方法
JP7436964B1 (ja) 2022-07-29 2024-02-22 住友電気工業株式会社 鋼線、及び鋼線の製造方法

Also Published As

Publication number Publication date
JP4719320B2 (ja) 2011-07-06
EP2447382A1 (en) 2012-05-02
EP2447382A4 (en) 2017-08-02
EP2447382B1 (en) 2019-03-06
JPWO2010150450A1 (ja) 2012-12-06
TWI412608B (zh) 2013-10-21
TW201102443A (en) 2011-01-16
CN102066599A (zh) 2011-05-18
US20110168302A1 (en) 2011-07-14
CN102066599B (zh) 2013-11-06
KR101260598B1 (ko) 2013-05-03
KR20110020256A (ko) 2011-03-02

Similar Documents

Publication Publication Date Title
JP4719320B2 (ja) 高強度極細鋼線及びその製造方法
JP4943564B2 (ja) ソーワイヤ用素線及びその製造方法
JP5157230B2 (ja) 伸線加工性の優れた高炭素鋼線材
JP6264462B2 (ja) 伸線加工用鋼線
JP3954338B2 (ja) 耐ひずみ時効脆化特性および耐縦割れ性に優れる高強度鋼線およびその製造方法
WO2016002413A1 (ja) 鋼線用線材および鋼線
CA2980886C (en) High-carbon steel wire material with excellent wire drawability, and steel wire
WO2015119241A1 (ja) フィラメント
WO2018117157A1 (ja) 線材
JP2014055316A (ja) 高強度鋼線用線材
JP3777166B2 (ja) 高強度極細鋼線の製造方法
JP4464511B2 (ja) 延性及び疲労特性の優れた高強度極細鋼線の製造方法
JP4555711B2 (ja) 延性に優れた高強度極細鋼線
JP6724400B2 (ja) 強度と延性のバランスに優れた高強度極細鋼線及びその製造方法
JP2010275626A (ja) 耐断線性に優れた高強度極細鋼線及びその製造方法
JP6922726B2 (ja) 熱間圧延線材
JP2002212676A (ja) ワイヤソー用鋼線およびその製造方法
JP3641056B2 (ja) 高強度極細鋼線
WO2024024401A1 (ja) 鋼線、及び鋼線の製造方法
JP6946891B2 (ja) 高強度鋼線
JPH06346190A (ja) 疲労特性に優れた極細鋼線

Legal Events

Date Code Title Description
WWE Wipo information: entry into national phase

Ref document number: 201080001860.7

Country of ref document: CN

ENP Entry into the national phase

Ref document number: 2010540974

Country of ref document: JP

Kind code of ref document: A

WWE Wipo information: entry into national phase

Ref document number: 12994169

Country of ref document: US

WWE Wipo information: entry into national phase

Ref document number: 2010778818

Country of ref document: EP

ENP Entry into the national phase

Ref document number: 20107028271

Country of ref document: KR

Kind code of ref document: A

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

Ref document number: 10778818

Country of ref document: EP

Kind code of ref document: A1

NENP Non-entry into the national phase

Ref country code: DE