EP2546379B1 - Hochfester stahl und hochfester bolzen mit hervorragender beständigkeit gegen verzögerten bruch sowie herstellungsverfahren dafür - Google Patents

Hochfester stahl und hochfester bolzen mit hervorragender beständigkeit gegen verzögerten bruch sowie herstellungsverfahren dafür Download PDF

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EP2546379B1
EP2546379B1 EP20110753527 EP11753527A EP2546379B1 EP 2546379 B1 EP2546379 B1 EP 2546379B1 EP 20110753527 EP20110753527 EP 20110753527 EP 11753527 A EP11753527 A EP 11753527A EP 2546379 B1 EP2546379 B1 EP 2546379B1
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Prior art keywords
steel material
less
bolt
high strength
mass
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French (fr)
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EP2546379A4 (de
EP2546379A1 (de
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Daisuke Hirakami
Tetsushi Chida
Toshimi Tarui
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Nippon Steel Corp
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Nippon Steel and Sumitomo Metal Corp
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—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
    • 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
    • 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
    • C21D9/0093—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for screws; for bolts
    • 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
    • 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
    • C22C38/002—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
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    • 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
    • C22C38/18—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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of 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
    • C22C38/18—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
    • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel with boron
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C8/00—Solid state diffusion of only non-metal elements into metallic material surfaces; Chemical surface treatment of metallic material by reaction of the surface with a reactive gas, leaving reaction products of surface material in the coating, e.g. conversion coatings, passivation of metals
    • C23C8/06—Solid state diffusion of only non-metal elements into metallic material surfaces; Chemical surface treatment of metallic material by reaction of the surface with a reactive gas, leaving reaction products of surface material in the coating, e.g. conversion coatings, passivation of metals using gases
    • C23C8/08—Solid state diffusion of only non-metal elements into metallic material surfaces; Chemical surface treatment of metallic material by reaction of the surface with a reactive gas, leaving reaction products of surface material in the coating, e.g. conversion coatings, passivation of metals using gases only one element being applied
    • C23C8/24—Nitriding
    • C23C8/26—Nitriding of ferrous surfaces
    • 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
    • C21D2211/008—Martensite
    • 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
    • C21D2221/00—Treating localised areas of an article
    • C21D2221/10—Differential treatment of inner with respect to outer regions, e.g. core and periphery, respectively

Definitions

  • the present invention relates to a high strength steel material which is used for wire rods, PC steel bars (steel bars for prestressed concrete use), etc., more particularly relates to a high strength steel material and high strength bolts of a tensile strength of 1300 MPa or more which are excellent in delayed fracture resistance and methods for the production of the same.
  • the high strength steel which is used in large amounts for machines, automobiles, bridges, and building structures is medium carbon steel with an amount of C of 0.20 to 0.35%, for example, SCr, SCM, etc. defined by JIS G 4104 and JIS G 4105 which is quenched and tempered.
  • C 0.20 to 0.35%
  • the method of making the steel structure a bainite structure or the method of refining the prior austenite grains is effective.
  • PLT 1 discloses steel which is refined in prior austenite grains and improved in delayed fracture resistance
  • PLT's 2 and 3 disclose steels which suppress segregation of steel ingredients to improve the delayed fracture resistance.
  • refinement of prior austenite grains or with suppression of segregation of ingredients it is difficult to greatly improve the delayed fracture resistance.
  • a bainite structure contributes to improvement of the delayed fracture resistance, but formation of a bainite structure requires suitable additive elements or heat treatment, so the cost of the steel rises.
  • PLT's 4 to 6 disclose wire rods for high strength bolts containing 0.5 to 1.0 mass% of C in which an area ratio 80% or more of the pearlite structure is strongly drawn to impart 1200N/mm 2 or more strength and excellent delayed fracture resistance.
  • the wire rods which are described in PLT's 4 to 6 are high in cost due to the drawing process. Further, manufacture of thick wire rods is difficult.
  • PLT 7 discloses a coil spring using oil-tempered wire of a hardness of the inside cross-section of Hv550 or more to suppress the occurrence of delayed fracture after cold coiling.
  • the surface hardness of the product after nitriding is Hv900 or more. Under high load stress such as with bolts or PC steel bars, the delayed fracture characteristics are low. If the corrosive environment becomes severer, there is the problem of delayed fracture occurring.
  • Both PLT 8 and PLT 9 disclose a steel material which is excellent in delayed fracture resistance composed of steel of a required composition which is nitrided and mainly consists of a tempered martensite structure.
  • the high strength steel material which is disclosed in PLT 8 and PLT 9 exhibit delayed fracture resistance even under a corrosive environment containing hydrogen.
  • the present invention in view of this current situation, has as its object to provide a high strength steel material (wire rod or PC steel bar) and high strength bolt which exhibit excellent delayed fracture resistance even under a severe corrosive environment and methods of production for producing these inexpensively.
  • the inventors engaged in intensive research on the techniques for solving the above problem. As a result, they learned that if (a) introducing into the steel material predetermined amounts of V and/or Mo - which form precipitates which trap hydrogen, (b) decarburizing and nitriding the surface of the steel material (b1) to form a low carbon region to suppress hardening and (b2) to form a nitrided layer to obstruct absorption of hydrogen, the delayed fracture resistance is remarkably improved.
  • the present invention was made based on the above discovery and has as its gist the following:
  • FIG. 1(a) schematically shows a hydrogen evolution rate curve obtained by hydrogen analysis by the thermal desorption analysis. As shown in FIG. 1(a) , the amount of release of diffusible hydrogen reaches a peak near 100°C.
  • a sample is raised in temperature by 100°C/h and the cumulative value of the amount of hydrogen which is released from room temperature to 400°C is defined as the amount of diffusible hydrogen.
  • the amount of released hydrogen can be measured by a gas chromatograph.
  • the minimum amount of diffusible hydrogen at which delayed fracture occurs is referred to as the "critical diffusible hydrogen content”.
  • the limit amount of diffusible hydrogen differs according to the type of the steel.
  • FIG. 1(b) schematically shows the relationship between the fracture time obtained by a constant load delayed fracture test of the steel material and the amount of diffusible hydrogen. As shown in FIG. 1(b) , if the amount of diffusible hydrogen is great, the fracture time is short, while if the amount of diffusible hydrogen is small, the fracture time is long.
  • the inventors studied the composition of the steel material predicated on trapping and rendering harmless the hydrogen due to ingress in the steel material by fine precipitates so as to increase the critical diffusible hydrogen content.
  • the inventors confirmed that if including suitable amounts of one or both of V and Mo and causing fine precipitates composed of carbides, nitrides, and/or carbonitrides of V or Mo to form in the steel material, it is possible to increase the critical diffusible hydrogen content.
  • the hexagonal carbides, nitrides, and carbonitrides of Mo have a larger effect of increase of the critical diffusible hydrogen content than the effect of increase of the amount of absorption of hydrogen, so contribute to the improvement of the delayed fracture resistance, but compared with the NaCl-type structures of carbides, nitrides, and carbonitrides of V, the effect of improvement of strength by fine precipitation is small.
  • NaCl-type structures of carbides, nitrides, and carbonitrides of V are excellent in the effect of improvement of strength due to fine precipitation.
  • the NaCl-type structures of carbides, nitrides, and carbonitrides of V increase the critical diffusible hydrogen content, while also increase the amount of absorption of hydrogen, so the effect of improvement of the delayed fracture resistance is smaller compared with the hexagonal carbides, nitrides, and carbonitrides of Mo.
  • the inventors engaged in in-depth studies to suitably set the contents of Mo and V. As a result, they learned that if including one or both of V: 1.5 mass% or less and Mo: 3.0 mass% or less and making the contents of V and Mo (mass%) satisfy V+1/2Mo>0.4 mass%, the critical diffusible hydrogen content is increased and an excellent delayed fracture resistance is obtained.
  • the inventors discovered that by heating and rapid cooling at the time of nitriding, compressive residual stress occurs at the steel material surface and the delayed fracture resistance is improved.
  • the delayed fracture resistance is improved.
  • the nitrogen concentration becomes higher, so the delayed fracture resistance is remarkably improved.
  • the high strength steel material and high strength bolt of the present invention are composed of predetermined compositions of ingredients and have a nitrided layer and a low carbon region simultaneously present on the surface.
  • the high strength steel material and high strength bolt of the present invention there is a region with a nitrogen concentration of 12.0 mass% or less and higher than the nitrogen concentration of the steel material by 0.02 mass% or more and with a carbon concentration of 0.05 mass% or more and 0.9 time or less the steel material (low carbon nitrided layer).
  • the carbon concentration at the location deeper than the low carbon region is equal to the carbon concentration of the steel material and the nitrogen concentration is higher than the nitrogen concentration of the steel material.
  • the result is a low carbon region with a carbon concentration of 0.05 mass% or more and 0.9 times or less of the carbon concentration of the steel material and with contents of other elements equal to the steel material is present under the nitrided layer.
  • the low carbon region is a region with a carbon concentration of 0.05 mass% or more and 0.9 time or less the carbon concentration of the high strength steel material or high strength bolt.
  • a low carbon region is formed at a depth of 100 ⁇ m or more to 1000 ⁇ m from the steel material surface.
  • the depth and carbon concentration of the low carbon region adjust the heating atmosphere, heating temperature, and holding time at the time of heat treatment which forms the low carbon region. For example, if the carbon potential of the heating atmosphere is low, the heating temperature is high, and the holding time is long, the low carbon region becomes deeper and the carbon concentration of the low carbon region falls.
  • the carbon concentration of the low carbon region is less than 0.05 mass%, this becomes less than half of the lower limit 0.10 mass% of the carbon concentration of the steel material, so it is not possible to secure a predetermined strength and hardness by the low carbon region. If the carbon concentration of the low carbon region is over 0.9 time the carbon concentration of the steel material, this is substantially equal to the carbon concentration of the steel material and the effect of presence of the low carbon region ends up becoming weaker.
  • the low carbon region was defined as a region where the carbon concentration is 0.05 mass% or more and 0.9 time or less the carbon concentration of the steel material.
  • the carbon concentration of the low carbon region is 0.05 mass% or more and 0.9 time or less of the carbon concentration of the steel material, it is possible to reduce the amount of increase in the surface hardness due to formation of the nitrided layer. As a result, the hardness of the surface of the steel material becomes equal to the hardness of the steel material or lower than the hardness of the steel material and can prevent a reduction of the critical diffusible hydrogen content.
  • the depth (thickness) of the low carbon region was made a depth (thickness) of 100 ⁇ m or more from the surface of the steel material or bolt so that the effect is obtained.
  • the depth (thickness) of the low carbon region is preferably greater in depth (thickness), but if over 1000 ⁇ m, the strength of the steel material as a whole or the bolt as a whole falls, so the depth (thickness) of the low carbon region is given an upper limit of 1000 ⁇ m.
  • the nitrided layer is a region with a nitrogen concentration of 12.0 mass% or less and higher than the nitrogen concentration of the steel material or bolt by 0.02 mass% or more. Further, the nitrided layer is formed by a thickness of 200 ⁇ m or more from the surface of the steel material or bolt.
  • the thickness and nitrogen concentration of the nitrided layer can be adjusted by the heating atmosphere, heating temperature, and holding time at the time of nitriding. For example, if the concentration of ammonia or nitrogen in the heating atmosphere is high, the heating temperature is high, and the holding time is long, the nitrided layer becomes thicker and the nitrogen concentration of the nitrided layer becomes higher.
  • the nitrogen concentration of the nitrided layer is higher than the nitrogen concentration of the steel material, it is possible to reduce the amount of absorption of hydrogen in the steel material from a corrosive environment, but if the difference of the nitrogen concentration of the nitrided layer and the nitrogen concentration of the steel material is less than 0.02 mass%, the effect of reduction of the amount of absorption of hydrogen cannot be sufficiently obtained. For this reason, the nitrogen concentration of the nitrided layer was made a concentration higher than the nitrogen concentration of the steel material by 0.02 mass% or more.
  • the steel material surface is formed with a nitrided layer which has a nitrogen concentration of 12.0 mass% or less and higher than the nitrogen concentration of the steel material by 0.02 mass% or more and a depth of 200 ⁇ m or more from the surface, the amount of absorption of hydrogen in the steel material from the corrosive environment is greatly reduced.
  • the nitrided layer was limited to a thickness (depth) of 200 ⁇ m or more from the surface of the steel material or bolt so that the effect is obtained.
  • the upper limit of the thickness of the nitrided layer is not particularly defined, but if the thickness is over 1000 ⁇ m, the productivity falls and a rise in cost is invited, so 1000 ⁇ m or less is preferable.
  • the depth (thickness) of the low carbon region which is formed on the high strength steel material or high strength bolt of the present invention can be found from the curve of the carbon concentration from the surface of the steel material or bolt.
  • a cross-section of a steel material or bolt which has a low carbon region and nitrided layer on the surface is polished and an Energy Dispersive X-ray Spectroscopyter (below, sometimes referred to as "EDX”) or a Wavelength Dispersive X-ray Spectroscopy (below, sometimes referred to as "WDS”) is used for line analysis to measure the carbon concentration in a depth direction from the surface.
  • EDX Energy Dispersive X-ray Spectroscopyter
  • WDS Wavelength Dispersive X-ray Spectroscopy
  • FIG. 2 shows the method of finding the depth (thickness) of the low carbon region from the curve of the carbon concentration which is obtained by EDX. That is, FIG. 2 is a view which shows the relationship between the distance from the steel material surface, obtained by measuring the carbon concentration in the depth direction from the surface using EDX, and the carbon concentration.
  • the carbon concentration increases along with the increased distance (depth) from the steel material surface. This is because due to decarburization, a low carbon region is formed on the surface of the steel material.
  • the carbon concentration is substantially constant (average carbon concentration "a").
  • the average carbon concentration "a” is the carbon concentration of the region not affected by the decarburization and is equal to the amount of carbon of the steel material before decarburization.
  • the chemical analysis value of the carbon concentration of the steel material is made the reference value when finding the depth of the low carbon region.
  • the thickness (depth) of the nitrided layer can be found from the change of the nitrogen concentration from the surface of the steel material or bolt in the same way as the low carbon region. Specifically, a cross-section of the steel material or bolt which has a low carbon region and nitrided layer on the surface is polished and an EDX or WDS is used for line analysis to measure the nitrogen concentration in the depth direction from the surface.
  • FIG. 3 shows the method of finding the thickness (depth) of the nitrided layer from the nitrogen concentration curve obtained by an Energy Dispersion type X-ray Spectrometer (EDX). That is, FIG. 3 is a view showing the relationship between the distance from the steel material surface and the nitrogen concentration which is obtained by measuring the nitrogen concentration in the depth direction from the surface using EDX.
  • EDX Energy Dispersion type X-ray Spectrometer
  • the nitrogen concentration decreases, but in the region not affected by nitriding, the carbon concentration is substantially constant (average nitrogen concentration).
  • the average nitrogen concentration is a range of nitrogen concentration not affected by nitriding and is equal to the amount of nitrogen of the steel material before nitriding. Therefore, in the present invention, the chemical analysis value of the nitrogen concentration of the steel material is made the reference value when finding the thickness of the nitrided layer.
  • the depth of the low carbon region and the thickness of the nitrided layer are found by obtaining simple averages of the values which were measured at any five locations at the cross-section of the steel material or bolt.
  • the carbon concentration and nitrogen concentration of the steel material may be found by measuring the carbon concentration and nitrogen concentration at a position sufficiently deeper than the depth of the low carbon region and nitrided layer, for example, a position at a depth of 2000 ⁇ m or more from the surface. Further, it is also possible to obtain an analytical sample from a position at a depth of 2000 ⁇ m or more from the surface of the steel material or bolt and chemically analyze it to find them.
  • the delayed fracture is remarkably improved by the synergistic effect of (1) suppression of the amount of absorption of hydrogen due to the formation of a nitrided layer at the low carbon region which is formed at the steel material surface and (2) increase of the critical diffusible hydrogen content due to the formation of the low carbon region at the steel material surface.
  • the surface of the steel material has a nitrided layer and a low carbon region copresent on it, whereby the amount of absorption of hydrogen in the steel material can be suppressed to 0.5 ppm or less and the critical diffusible hydrogen content of the steel material can be raised to 2.00 ppm or more.
  • the % according to the composition mean mass%.
  • C is an essential element in securing the strength of a steel material. If less than 0.10%, the required strength is not obtained, while if over 0.55%, the ductility and toughness fall and the delayed fracture resistance also falls, so the content of C was made 0.10 to 0.55%.
  • Si is an element which improves strength by solution strengthening. If less than 0.01%, the effect of addition is insufficient, while if over 3%, the effect becomes saturated, so the content of Si was made 0.01 to 3%.
  • Mn is an element which not only performs deoxidation and desulfurization, but also gives a martensite structure, so lowers the transformation temperature of the pearlite structure or bainite structure to raise the hardenability. If less than 0.1%, the effect of addition is insufficient, while if over 2%, it segregates at the grain boundary at the time of heating of austenite to embrittle the grain boundary and degrades the delayed fracture resistance, so the content of Mn was made 0.1 to 2%.
  • the high strength steel material or high strength bolt of the present invention contains one or both of V and Mo to improve the delayed fracture resistance.
  • V is an element which causes the formation of fine precipitates of carbides, nitrides, and/or carbonitrides in the steel material. These fine precipitates act to trap and render harmless the hydrogen due to ingress in the steel material and increase the critical diffusible hydrogen content. Further, the fine precipitates contribute to the improvement of strength of the steel material.
  • V is an element which lowers the transformation temperature of the pearlite structure or bainite structure and increases the hardenability and also raises the resistance to softening during tempering at the time of tempering to contribute to improvement of the strength.
  • the content of V exceeds 1.5%, the solution temperature for precipitation strengthening becomes higher and, further, the ability for trapping hydrogen also becomes saturated, so the content was made 1.5% or less. If the content of V is less than 0.05%, an effect of improvement of strength due to the formation of fine precipitates is not sufficiently obtained, so 0.05% or more is preferable.
  • Mo is an element which causes the formation of fine precipitates of carbides, nitrides, and/or carbonitrides in the steel material.
  • the fine precipitates act to trap and render harmless the hydrogen due to ingress in the steel material and increase the critical diffusible hydrogen content. Further, fine precipitates contribute to an improvement of strength of the steel material and increases the critical diffusible hydrogen content to contribute to the improvement of the delayed fracture resistance.
  • Mo is an element which lowers the transformation temperature of the pearlite structure or bainite structure to raise the hardenability and raises the resistance to softening during tempering at the time of tempering to contribute to the improvement of the strength.
  • the content of Mo exceeds 3.0%, the solution temperature for precipitation strengthening becomes higher and, further, the ability for trapping hydrogen also becomes saturated, so the content was made 3.0% or less. If the content of Mo is less than 0.4%, an effect of increase of the critical diffusible hydrogen content due to the formation of fine precipitates is not sufficiently obtained, so 0.4% or more is preferable.
  • V+1/2Mo In the present invention, the contents of Mo and V have to satisfy V+1/2Mo>0.4%. If the contents of V and Mo satisfy this formula, the amount of V whereby not only the critical diffusible hydrogen content, but also the amount of absorption of hydrogen increases becomes relatively smaller than the amount of Mo, so the critical diffusible hydrogen content increases and an excellent delayed fracture resistance is obtained.
  • the high strength steel material or high strength bolt of the present invention may further contain one or more of Cr, Nb, Cu, Ni, and B in a range not impairing the excellent delayed fracture resistance for the purpose of improving the strength.
  • Cr is an element which lowers the transformation temperature of the pearlite structure or bainite structure to raise the hardenability and, further, raises the resistance to softening during tempering at the time of tempering to contribute to the improvement of the strength. If less than 0.05%, the effect of addition is not sufficiently obtained, while if over 1.5%, deterioration of the toughness is invited, so the content of Cr was made 0.05 to 1.5%.
  • Nb is an element which raises the hardenability and the tempering softening resistance to contribute to the improvement of the strength. If less than 0.001%, the effect of addition is not sufficiently obtained. If over 0.05%, the effect of addition becomes saturated, so the content of Nb was made 0.001 to 0.05%.
  • Cu is an element which contributes to the improvement of the hardenability, increase of the temper softening resistance, and improvement of strength by the precipitation effect. If less than 0.01%, the effect of addition is not sufficiently obtained, while if over 4%, grain boundary embrittlement occurs and the delayed fracture resistance deteriorates, so the content of Cu was made 0.01 to 4%.
  • Ni is an element which raises the hardenability and is effective for improvement of the ductility and toughness which fall along with increased strength. If less than 0.01%, the effect of addition is not sufficiently obtained, while if over 4%, the effect of addition becomes saturated, so the content of Ni was made 0.01 to 4%.
  • B is an element which suppresses grain boundary fracture and is effective for improvement of the delayed fracture resistance. Furthermore, B is an element which segregates at the austenite grain boundary and remarkably raises the hardenability. If less than 0.0001%, the effect of addition cannot be sufficiently obtained, while if over 0.005%, B carbides and Fe borocarbides form at the grain boundaries, grain boundary embrittlement occurs, and delayed fracture resistance falls, so the content of B is made 0.0001 to 0.005%.
  • the high strength steel material and high strength bolt of the present invention may further contain, for the purpose of refining the structure, one or more of Al, Ti, Mg, Ca, and Zr in a range not detracting from the excellent delayed fracture resistance.
  • Al is an element which forms oxides or nitrides and prevents coarsening of austenite grains to suppress deterioration of the delayed fracture resistance. If less than 0.003%, the effect of addition is insufficient, while if over 0.1%, the effect of addition becomes saturated, so the content of Al is preferably 0.003 to 0.1%.
  • Ti also, like Al, is an element which forms oxides or nitrides to prevent coarsening of austenite grains and suppress deterioration of the delayed fracture resistance. If less than 0.003%, the effect of addition is insufficient, while if over 0.05%, the Ti carbonitrides coarsen at the time of rolling or working or at the time of heating in heat treatment and the toughness falls, so the content of Ti is preferably 0.003 to 0.05%.
  • Mg is an element which has a deoxidizing and desulfurizing effect and, further, forms Mg oxides, Mg sulfides, Mg-Al, Mg-Ti, and Mg-Al-Ti composite oxides or composite sulfides, etc. to prevent coarsening of austenite grains and suppress deterioration of delayed fracture resistance. If less than 0.0003%, the effect of addition is insufficient, while if over 0.01%, the effect of addition becomes saturated, so the content of Mg is preferably 0.0003 to 0.01%.
  • Ca is an element which has a deoxidizing and desulfurizing effect and, further, forms Ca oxides, Ca sulfides, Al, Ti, and Mg composite oxides or composite sulfides, etc. to prevent coarsening of austenite grains and suppress deterioration of delayed fracture resistance. If less than 0.0003%, the effect of addition is insufficient, while if over 0.01%, the effect of addition becomes saturated, so the content of Ca is preferably 0.0003 to 0.01%.
  • Zr is an element which forms Zr oxides, Zr sulfides, Al, Ti, Mg, and Zr composite oxides or composite sulfides, etc. to prevent coarsening of austenite grains and suppress deterioration of delayed fracture resistance. If less than 0.0003%, the effect of addition is insufficient, while if over 0.01%, the effect of addition becomes saturated, so the content of Zr is preferably 0.0003 to 0.01%.
  • the steel structure of the present invention is mainly tempered martensite, so the structure is excellent in ductility and toughness even if the tensile strength is 1300 MPa or more.
  • the steel structure of the present invention is preferably a structure where the area ratio of the tempered martensite in the region excluding the low carbon region and nitrided layer is 85% or more and the balance is composed of one or more of residual austenite, bainite, pearlite, and ferrite.
  • the area ratio of the tempered martensite is measured at a deeper position between the depth at which the carbon concentration becomes constant in the carbon concentration curve which is shown in FIG. 2 and the depth where the nitrogen concentration becomes constant in the nitrogen concentration curve which is shown in FIG. 3 .
  • the area ratio of martensite can be found by observing the cross-section of the steel material using an optical microscope and measuring the area of martensite per unit area. Specifically, the cross-section of the steel material is etched by a Nital etching solution, the areas of martensite in five fields in a range of 0.04 mm 2 are measured, and the average value is calculated.
  • compressive residual stress of the steel material surface occurs due to the heating and rapid cooling at the time of nitriding whereby the delayed fracture resistance is improved. If the compressive residual stress occurs by 200 MPa or more, the delayed fracture resistance is improved, so the compressive residual stress of the surface of the steel material of the present invention is preferably 200 MPa or more.
  • the compressive residual stress can be measured by X-ray diffraction. Specifically, the residual stress of the steel material surface is measured, then the steel material surface is etched 25 ⁇ m at a time by electrolytic polishing and the residual stress in the depth direction is measured. It is preferable to measure any three locations and use the average value of the same.
  • the tensile strength becomes 1300 MPa or more, the frequency of occurrence of delayed fracture remarkably increases. Therefore, if the tensile strength is 1300 MPa or more, the delayed fracture resistance of the steel material of the present invention on which a low carbon region and nitrided layer are formed on the surface is remarkably excellent.
  • the upper limit of the tensile strength of the present invention is not particularly limited, but over 2200 MPa is technically difficult at the present point of time, so 2200 MPa is provisionally made the upper limit.
  • the tensile strength may be measured based on JIS Z 2241.
  • the method of production of a steel material of the present invention is composed of a decarburization step of heating a steel material of a required composition (wire rod or PC steel bar or steel material worked to a predetermined shape) to decarburize it, a hardening step of cooling the decarburized steel material to make the steel structure a mainly martensite structure, and a step of nitriding the hardened steel material at over 500°C to 650°C or less.
  • the structure of the steel material of the present invention becomes a structure of mainly tempered martensite.
  • the steel material of the present invention is decarburized to make the carbon concentration, down from the surface of the steel material by a depth of 100 ⁇ m or more to 1000 ⁇ m or less, 0.05% or more and 0.9 time or less the carbon concentration of the steel material.
  • the atmosphere in the heating furnace is, for example, adjusted to a concentration of methane gas to make it weakly decarburizing and form a low carbon region.
  • the heating temperature in the decarburization is preferably Ac 3 to 950°C. By heating to Ac 3 or more, it is possible to make the steel structure austenite, promote decarburization from the surface layer, and easily form a low carbon region.
  • the upper limit of the heating temperature is preferably 950°C in the point that this suppresses coarsening of the crystal grains and improves the delayed fracture resistance.
  • the holding time at the heating temperature is preferably 30 to 90 minutes. By holding at the heating temperature for 30 minutes or more, it is possible to sufficiently secure the depth of the low carbon region and possible to make the steel structure uniform. If considering the productivity, the holding time at the heating temperature is preferably 90 minutes or less
  • the heated steel material is cooled to obtain a mainly martensite structure.
  • the heated steel material may be oil quenched as it is for hardening.
  • the area ratio of the tempered martensite is preferably 85% or more, so the area ratio of the martensite after hardening is preferably 85% or more.
  • the cooling rate in the range from 700 to 300°C 5°C/s or more. if the cooling rate is less than 5°C/s, sometimes the area ratio of the martensite becomes less than 85%.
  • a steel material with a steel structure of mainly martensite and formed with a low carbon region at the surface layer is nitrided. Due to the nitriding, a nitrided layer is formed with a thickness from the steel material surface of 200 ⁇ m or more and a nitrogen concentration of 12.0% or less and higher than the nitrogen concentration of the steel material by 0.02% or more.
  • the steel material is tempered to make the steel structure a mainly tempered martensite structure and to form fine precipitates which trap hydrogen.
  • the steel material contains one or both of V and Mo, so in the nitriding step, fine carbides, nitrides, and/or carbonitrides (fine precipitates) which act to trap the hydrogen are formed.
  • the nitriding is performed by, for example, heating the steel material in an atmosphere containing ammonia or nitrogen.
  • the nitriding is preferably performed by holding the sample at over 500°C to 650°C or less for 30 to 90 minutes. If the nitriding temperature exceeds 650°C, the steel material falls in strength, so the nitriding temperature is made 650°C or less.
  • the nitriding temperature is 500°C or less, fine precipitates do not sufficiently form, so the nitriding temperature is made over 500°C. Further, if the nitriding temperature is made over 500°C, the time required for diffusion of nitrogen from the steel material surface becomes shorter, the treatment time is shortened, and the productivity rises.
  • the nitriding time is less than 30 minutes, the depth of the nitrided layer is liable not to reach a depth of 200 ⁇ m or more from the surface, so the nitriding time is preferably 30 minutes or more.
  • the upper limit of the nitriding time is not defined, but in the present invention, the nitriding temperature is high, so even at 90 minutes or less, a sufficient thickness of a nitrided layer can be formed.
  • the gas nitriding method nitrocarburizing method, plasma nitriding method, salt bath nitriding method, or other general nitriding method may be used.
  • the method of production of the bolt of the present invention is composed of a working step of working the steel material of the present invention having the required composition into a bolt, a decarburization step of heating the bolt to decarburize it, a hardening step of cooling the heated bolt to make the steel structure a mainly martensite structure, and a nitriding step of nitriding the hardened bolt at a temperature of over 500°C to 650°C or less.
  • the steel structure of the bolt becomes a mainly tempered martensite structure.
  • the steel material wire rod is cold forged and rolled to form a bolt.
  • the method of production of the bolt of the present invention differs from the method of production of the steel material of the present invention only in the working step for working the steel material into a bolt shape, so the explanation of the other steps will be omitted.
  • the method of production of the steel material of the present invention and the method of production of the bolt of the present invention preferably performs rapid cooling, after nitriding, in a range from 500 to 200°C by a cooling rate of 10 to 100°C/s.
  • rapid cooling after nitriding, it is possible to make the compressive residual stress of the surface of the steel material or bolt 200 MPa or more. Due to the presence of this compressive residual stress, the delayed fracture resistance is improved more.
  • Molten steels of the compositions of ingredients which are shown in Table 1 were cast in accordance with an ordinary method.
  • the cast slabs were hot worked to obtain steel materials (wire rods).
  • the steel materials were heated to Ac 3 to 950°C and cooled as is for hardening.
  • the atmosphere in the heating furnace was controlled to be weakly decarburizing.
  • the hardening was performed by oil quenching so that the cooling rate in the range of 700 to 300°C became 5°C/s or more.
  • the depth of the low carbon region was investigated by the carbon potential of the atmosphere of the heating furnace, heating temperature, and holding time.
  • Table 1 Steel type Composition (mass%) C Si Mn Cr V Mo Nb Cu Ni B Al Ti Mg Ca Zr Fe V+1/2Mo A1 0.11 2.2C 1.52 0.53 0.21 0.98 0.049. 3.85 2.25 0.0045 0.032 - - 0.0003 - Bal.
  • the steel materials were nitrided by nitrocarburizing to form nitrided layers. After nitriding, the materials were rapidly cooled in the range of 500 to 200°C by the cooling rates which are shown in Table 2 (cooling rates after tempering) to obtain the high strength steel materials of Manufacturing Nos. 1 to 25.
  • the nitriding was performed at the temperatures which are shown in Table 2 while making the ammonia volume ratios in the treatment gas atmosphere 30 to 50% and making the treatment times 30 to 90 minutes.
  • the nitrided layers were adjusted in thickness by changing the heating temperature and the holding time.
  • the nitrided layers were adjusted in nitrogen concentration by changing the ammonia volume ratio in the treatment gas atmosphere.
  • Table 2 Man. No. Steel type Tempered martensite ratio (%) Strength (MPa) Low carbon region depth ( ⁇ m) Nitriding Nitrided layer thickness ( ⁇ m) Compressive residual stress (MPa) Amount of absorption of hydrogen (ppm) Critical diffusible hydrogen content of delayed fracture (ppm) Delayed fracture presence Remarks Temperature (°C) Cooling rate (°C/s) 1 A1 90 1312 130 580 35 365 306 0.23 3.50 No Inv. ex.
  • the tempered martensite ratios, tensile strengths, low carbon region depths, nitrided layer thicknesses, compressive residual stresses, amounts of absorption of hydrogen, limit amounts of diffusible hydrogen, and delayed fracture resistances of the high strength steel materials of Manufacturing Nos. 1 to 25 (Table 2) and the high strength bolts of Manufacturing Nos. 26 to 40 (Table 3) were measured by the methods which are shown below. The results are shown in Table 2 and Table 3 together.
  • the tempered martensite ratio was found by polishing the cross-section of each of the high strength steel materials of Manufacturing Nos. 1 to 25 and the high strength bolts of Manufacturing Nos. 26 to 40, etching by a Nital etching solution, using an optical microscope to measure the areas of the martensite in five fields in a 0.04 mm 2 range, and finding the average value.
  • the structure of the remaining part of the tempered martensite was a balance of one or more of austenite, bainite, pearlite, and ferrite.
  • the tensile strength was measured based on JIS Z 2241.
  • a cross-section of each of the high strength steel materials of Manufacturing Nos. 1 to 25 and the high strength bolts of Manufacturing Nos. 26 to 40 was polished and measured for the carbon concentration and nitrogen concentration in the depth direction from the surface using an EDX at any five locations in the longitudinal direction.
  • the depth (thickness) of the region where the carbon concentration is 0.9 time or less of the carbon concentration of the steel material was defined as the "low carbon region depth"
  • the depth (thickness) of the region where the nitrogen concentration is higher than the nitrogen concentration of the steel material by 0.02% or more was defined as the "nitrided layer thickness”.
  • the low carbon region depth and the nitrided layer thickness were the averages of values measured at any five locations in the longitudinal direction.
  • An X-ray residual stress measurement apparatus was used to measure the compressive residual stress of the surface.
  • the residual stress of the surface of each of the high strength steel materials of Manufacturing Nos. 1 to 25 and the high strength bolts of Manufacturing Nos. 26 to 40 was measured, then the surface was etched by 25 ⁇ m at a time by electrolytic polishing and the residual stress in the depth direction was measured. Note that, the compressive residual stress was made the average of the values measured at any three locations.
  • a delayed fracture test piece of the shape which is shown in FIG. 4 was prepared and subjected to absorption of hydrogen.
  • the electrolytic hydrogen charge method was used to change the charge current and change the amount of absorption of hydrogen as shown by Table 2 and Table 3.
  • each delayed fracture test piece which was subjected to absorption of hydrogen was plated with Cd to prevent dissipation of the diffusible hydrogen.
  • the test piece was leaved at room temperature for 3 hours to even the concentration of hydrogen at the inside.
  • a delayed fracture test machine which is shown in FIG. 5 was used to run a constant load delayed fracture test applying a tensile load of 90% of the tensile strength to the test piece 1. Note that, in the test machine which is shown in FIG. 5 , when applying a tensile load to the test piece 1, a balance weight 2 was placed at one end of a lever having the fulcrum 3 as the fulcrum and the test piece 1 was placed at the other end to conduct the test.
  • the maximum value of the amount of diffusible hydrogen of a test piece 1 which did not fracture even after performing the constant load delayed fracture test for 100 hours or more was made the critical diffusible hydrogen content.
  • the amount of diffusible hydrogen of the test piece 1 was measured by raising the delayed fracture test piece in temperature at 100°C/h and measuring the cumulative value of the amounts of hydrogen which were released between room temperature to 400°C by a gas chromatograph.
  • delayed fracture When comparing the amount of absorption of hydrogen and the critical diffusible hydrogen content of delayed fracture and the critical diffusible hydrogen content is greater than the amount of absorption of hydrogen, delayed fracture does not occur. Conversely, if the critical diffusible hydrogen content is smaller than the amount of absorption of hydrogen, delayed fracture occurs.
  • the delayed fracture resistance was evaluated as "without delayed fracture” when the amount of absorption of hydrogen which is shown in Table 2 and Table 3 was less than critical diffusible hydrogen content and as "with delayed fracture” when the amount of absorption of hydrogen was the critical diffusible hydrogen content or more.
  • the amount of absorption of hydrogen was determined by preparing a test piece of each of the high strength steel materials of Manufacturing Nos. 1 to 25 and the high strength bolts of Manufacturing Nos. 26 to 40 and running an accelerated corrosion test of the pattern of temperature, humidity, and time which is shown in FIG. 6 for 30 cycles.
  • the corroded layer at the surface of the test piece was removed by sandblasting, then the hydrogen was analyzed by the Thermal desorption analysis. The amount of hydrogen which was released from room temperature to 400°C was measured to find the amount of absorption of hydrogen.
  • the high strength steel materials of Manufacturing Nos. 1 to 15 of the invention examples had a low carbon region depth of 100 ⁇ m or more and a nitrided layer thickness of 200 ⁇ m or more. All had a tensile strength of 1300 MPa or more, an amount of absorption of hydrogen of 0.5 ppm or less, a critical diffusible hydrogen content of 2.00 ppm or more, an amount of absorption of hydrogen of less than the critical diffusible hydrogen content, and no delayed fracture.
  • the high strength steel materials of Manufacturing Nos. 1 to 15 all had a tempered martensite rate of 50% or more and a structure of mainly tempered martensite. Further, the high strength steel materials of Manufacturing Nos. 1 to 15 all had a compressive residual stress of 200 MPa or more.
  • the high strength steel material of Manufacturing No. 16 of the comparative example was an example where the amount of C, the amount of Si, and the amount of Mn were small and the strength was low.
  • Manufacturing No. 17 is an example where the amount of C was large
  • Manufacturing No. 18 is an example where the amount of Mn was large
  • Manufacturing No. 19 is an example where the amount of Cr was large
  • Manufacturing No. 21 is an example where the amount of Cu was large
  • Manufacturing No. 22 is an example where the amount of B was large
  • Manufacturing No. 20 is an example where V+1/2Mo was low. In each case, the critical diffusible hydrogen content was low and the resistance was "with delayed fracture".
  • Manufacturing No. 23 is an example where the heating time of the hardening was short, the low carbon region depth was less than 100 ⁇ m, the critical diffusible hydrogen content was low, and the resistance was "with delayed fracture”.
  • Manufacturing No. 24 is an example where the nitriding time was short, the nitrided layer thickness was less than 200 ⁇ m, the amount of absorption of hydrogen was large, and the resistance was "with delayed fracture”.
  • Manufacturing No. 25 is an example in which the concentration of ammonia in the gas of the nitriding was lowered, so at a location down to the depth of 200 ⁇ m from the surface, the difference of the nitrogen concentration from the steel material became 0.01 mass%, the amount of absorption of hydrogen was larger, and the resistance was "with delayed fracture".
  • the high strength bolts of Manufacturing Nos. 26 to 40 of the invention examples had a low carbon region depth of 100 ⁇ m or more and a nitrided layer thickness of 200 ⁇ m or more. All had a tensile strength of 1300 MPa or more, an amount of absorption of hydrogen of 0.5 ppm or less, a critical diffusible hydrogen content of 2.00 ppm or more, an amount of absorption of hydrogen of less than the critical diffusible hydrogen content, and a resistance "without delayed fracture".
  • the high strength bolts of Manufacturing Nos. 26 to 40 all had a tempered martensite ratio of 50% or more, a structure of mainly tempered martensite, and a compressive residual stress of 200 MPa or more.
  • the present invention it is possible to provide a high strength steel material (wire rod or PC steel bar) and high strength bolt which exhibit excellent delayed fracture resistance even in a severe corrosive environment and a method of production enabling inexpensive production of these. Accordingly, the present invention is extremely high in applicability in industries manufacturing and using steel materials.

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Claims (14)

  1. Ein hochfestes Stahlmaterial, enthaltend
    in Massen%
    C: 0,10 bis 0,55 %,
    Si: 0,01 bis 3 % und
    Mn: 0,1 bis 2 %,
    weiter enthaltend eines oder beide aus
    V: 1,5 % oder weniger und Mo: 3,0 % oder weniger, wobei der Gehalt an V und Mo V+1/2Mo>0,4 % erfüllt,
    weiter enthaltend eines oder mehrere aus
    Cr: 0,05 bis 1,5 %,
    Nb: 0,001 bis 0,05 %,
    Cu: 0,01 bis 4 %,
    Ni: 0,01 bis 4 % und
    B: 0,0001 bis 0,005 %, weiter gegebenenfalls enthaltend eines oder mehrere aus Al: 0,003 bis 0,1 %,
    Ti: 0,003 bis 0,05 %,
    Mg: 0,0003 bis 0,41 %,
    Ca: 0,0003 bis 0,01 % und
    Zr: 0,0003 bis 0,01 % und
    mit einem Rest von Fe und unvermeidbaren Verunreinigungen, wobei die Struktur hauptsächlich eine getemperte Martensitstruktur ist,
    wobei die Oberfläche des Stahlmaterials versehen ist mit
    (a) einer Nitrierschicht mit einer Dicke von der Oberfläche des Stahlmaterials von 200 µm oder mehr und einer Stickstoffkonzentration von 12,0 Massen% oder weniger und um 0,02 Massen% oder mehr höher als die Stickstoffkonzentration des Stahlmaterials und
    (b) einer kohlenstoffarmen Region mit einer Tiefe von der Oberfläche des Stahlmaterials von 100 µm oder mehr bis 1000 µm oder weniger und mit einer Kohlenstoffkonzentration von 0,05 Massen% oder mehr und einer 0,9-fachen Kohlenstoffkonzentration des Stahlmaterials oder weniger und wobei der getemperte Martensit ein Flächenverhältnis von 85 % oder mehr aufweist.
  2. Ein hochfestes Stahlmaterial wie in Anspruch 1 dargelegt, dadurch gekennzeichnet, dass aufgrund des Vorhandenseins der Nitrierschicht und der kohlenstoffarmen Region die Menge der Absorption von Wasserstoff in das Stahlmaterial 0,5 ppm oder weniger beträgt und der kritische diffusionsfähige Wasserstoffgehalt des Stahlmaterials bei 2,00 ppm oder mehr liegt.
  3. Ein hochfestes Stahlmaterial wie in einem der Ansprüche 1 bis 2 dargelegt, dadurch gekennzeichnet, dass die Nitrierschicht eine Dicke von 1000 µm oder weniger aufweist.
  4. Ein hochfestes Stahlmaterial wie in einem der Ansprüche 1 bis 3 dargelegt, dadurch gekennzeichnet, dass das Stahlmaterial an der Oberfläche eine Druckeigenspannung von 200 MPa oder mehr aufweist.
  5. Ein hochfestes Stahlmaterial wie in einem der Ansprüche 1 bis 4 dargelegt, dadurch gekennzeichnet, dass das Stahlmaterial eine Zugfestigkeit von 1300 MPa oder mehr aufweist.
  6. Ein hochfester Bolzen, erhalten durch Bearbeiten eines Stahlmaterials, enthaltend in Massen%
    C: 0,10 bis 0,55 %,
    Si: 0,01 bis 3 % und
    Mn: 0,1 bis 2 %,
    weiter enthaltend eines oder beide aus
    V: 1,5 % oder weniger und
    Mo: 3,0 % oder weniger, wobei der Gehalt an V und Mo
    V+1/2Mo>0,4 % erfüllt,
    weiter enthaltend eines oder mehrere aus
    Cr: 0,05 bis 1,5%,
    Nb: 0,001 bis 0,05 %,
    Cu: 0,01 bis 4 %,
    Ni: 0,01 bis 4 % und
    B: 0,0001 bis 0,005 %, weiter gegebenenfalls enthaltend eines oder mehrere aus Al: 0,003 bis 0,1 %,
    Ti: 0,003 bis 0,05 %,
    Mg: 0,0003 bis 0,01 %,
    Ca: 0,0003 bis 0,01 % und
    Zr: 0,0003 bis 0,01 % und
    mit einem Rest von Fe und unvermeidbaren Verunreinigungen, wobei die Struktur eine hauptsächlich getemperte Martensitstruktur ist,
    wobei die Oberfläche des Bolzens versehen ist mit
    (a) einer Nitrierschicht mit einer Dicke von der Oberfläche des Bolzens von 200 µm oder mehr und einer Stickstoffkonzentration von 12,0 Massen% oder weniger und um 0,02 Massen% oder mehr höher als die Stickstoffkonzentration des Stahlmaterials und
    (b) einer kohlenstoffarmen Region mit einer Tiefe von der Oberfläche des Bolzens von 100 µm oder mehr bis 1000 µm oder weniger und mit einer Kohlenstoffkonzentration von 0,05 Massen% oder mehr und einer 0,9-fachen Kohlenstoffkonzentration des Stahlmaterials oder weniger und wobei das getemperte Martensit ein Flächenverhältnis von 85 % oder mehr aufweist.
  7. Ein hochfester Bolzen wie in Anspruch 6 dargelegt, dadurch gekennzeichnet, dass aufgrund des Vorhandenseins der Nitrierschicht und der kohlenstoffarmen Region die Menge der Absorption von Wasserstoff in den Bolzen 0,5 ppm oder weniger beträgt und der kritische diffusionsfähige Wasserstoffgehalt des Bolzens bei 2,00 ppm oder mehr liegt.
  8. Ein hochfester Bolzen wie in einem der Ansprüche 6 bis 7 dargelegt, dadurch gekennzeichnet, dass die Nitrierschicht eine Dicke von 1000 µm oder weniger aufweist.
  9. Ein hochfester Bolzen wie in einem der Ansprüche 6 bis 8 dargelegt, dadurch gekennzeichnet, dass der Bolzen an der Oberfläche eine Druckeigenspannung von 200 MPa oder mehr aufweist.
  10. Ein hochfester Bolzen wie in einem der Ansprüche 6 bis 9 dargelegt, dadurch gekennzeichnet, dass der Bolzen eine Zugfestigkeit von 1300 MPa oder mehr aufweist.
  11. Ein Verfahren zur Herstellung eines hochfesten Stahlmaterials wie in einem der Ansprüche 1 bis 5 dargelegt, gekennzeichnet durch
    (1) Erwärmen eines Stahlmaterials mit einer Zusammensetzung wie in einem der Ansprüche 1 oder 3 dargelegt, um eine kohlenstoffarme Region mit einer Tiefe von der Oberfläche des Stahlmaterials von 100 µm oder mehr bis 1000 µm oder weniger und mit einer Kohlenstoffkonzentration von 0,05 Massen% oder mehr und einer 0,9-fachen Kohlenstoffkonzentration des Stahlmaterials oder weniger zu bilden, anschließendes Abkühlen so wie es ist, um aus der Stahlmaterialstruktur hauptsächlich eine Martensitstruktur zu machen, dann
    (2) Nitrieren des Stahlmaterials bei über 500°C bis 650°C oder weniger, um auf der Oberfläche des Stahlmaterials eine Nitrierschicht mit einer Stickstoffkonzentration von 12,0 Massen% oder weniger und um 0,02 Massen% höher als die Stickstoffkonzentration des Stahlmaterials und mit einer Dicke von der Oberfläche des Stahlmaterials von 200 µm oder mehr zu bilden und um aus der Stahlmaterialstruktur hauptsächlich eine getemperte Martensitstruktur zu machen.
  12. Ein Verfahren zur Herstellung eines hochfesten Stahlmaterials wie in Anspruch 11 dargelegt, dadurch gekennzeichnet, dass die Nitrierschicht eine Dicke von 1000 µm oder weniger aufweist.
  13. Ein Verfahren zur Herstellung eines hochfesten Bolzens wie in einem der Ansprüche 6 bis 10 dargelegt,
    wobei das Verfahren zur Herstellung eines Bolzens mit hervorragender verzögerter Bruchbeständigkeit, gekennzeichnet ist durch
    (1) Erwärmen eines Bolzens, erhalten durch Bearbeiten eines Stahlmaterials mit einer Zusammensetzung wie in Anspruch 6 dargelegt, um eine kohlenstoffarme Region mit einer Tiefe von der Oberfläche des Bolzens von 100 µm oder mehr bis 1000 µm oder weniger und mit einer Kohlenstoffkonzentration von 0,05 Massen% oder mehr und mit einer 0,9-fachen Kohlenstoffkonzentration des Stahlmaterials oder weniger zu bilden, anschließendes Abkühlen so wie es ist, um aus der Stahlmaterialstruktur hauptsächlich eine Martensitestruktur zu machen, dann
    (2) Nitrieren des Bolzens bei über 500°C bis 650°C oder weniger, um auf der Oberfläche des Bolzens eine Nitrierschicht mit einer Stickstoffkonzentration von 12,0 Massen% oder weniger und um 0,02 Massen% höher als die Stickstoffkonzentration des Stahlmaterials und mit einer Dicke von der Oberfläche des Bolzens von 200 µm oder mehr zu bilden und um aus der Stahlmaterialstruktur hauptsächlich eine getemperte Martensitestruktur zu machen.
  14. Ein Verfahren zur Herstellung eines hochfesten Bolzens wie in Anspruch 13 dargelegt, dadurch gekennzeichnet, dass die Nitrierschicht eine Dicke von 1000 µm oder weniger aufweist.
EP20110753527 2010-03-11 2011-03-11 Hochfester stahl und hochfester bolzen mit hervorragender beständigkeit gegen verzögerten bruch sowie herstellungsverfahren dafür Not-in-force EP2546379B1 (de)

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