WO2013105344A1 - ボルト用鋼、ボルトおよびボルトの製造方法 - Google Patents
ボルト用鋼、ボルトおよびボルトの製造方法 Download PDFInfo
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- WO2013105344A1 WO2013105344A1 PCT/JP2012/080440 JP2012080440W WO2013105344A1 WO 2013105344 A1 WO2013105344 A1 WO 2013105344A1 JP 2012080440 W JP2012080440 W JP 2012080440W WO 2013105344 A1 WO2013105344 A1 WO 2013105344A1
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- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 79
- 239000010959 steel Substances 0.000 title claims abstract description 79
- 238000004519 manufacturing process Methods 0.000 title claims description 25
- XEEYBQQBJWHFJM-UHFFFAOYSA-N Iron Chemical compound [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 claims abstract description 41
- 229910052742 iron Inorganic materials 0.000 claims abstract description 9
- 239000012535 impurity Substances 0.000 claims abstract description 6
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- 238000010438 heat treatment Methods 0.000 claims description 22
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- 238000005096 rolling process Methods 0.000 claims description 12
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- 229910052760 oxygen Inorganic materials 0.000 claims description 9
- 229910052802 copper Inorganic materials 0.000 claims description 7
- 239000011261 inert gas Substances 0.000 claims description 5
- 229910052759 nickel Inorganic materials 0.000 claims description 3
- 239000005539 carbonized material Substances 0.000 claims description 2
- 239000001257 hydrogen Substances 0.000 abstract description 150
- 229910052739 hydrogen Inorganic materials 0.000 abstract description 150
- UFHFLCQGNIYNRP-UHFFFAOYSA-N Hydrogen Chemical compound [H][H] UFHFLCQGNIYNRP-UHFFFAOYSA-N 0.000 abstract description 137
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- 150000002431 hydrogen Chemical class 0.000 description 13
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- 229910021607 Silver chloride Inorganic materials 0.000 description 1
- UCKMPCXJQFINFW-UHFFFAOYSA-N Sulphide Chemical compound [S-2] UCKMPCXJQFINFW-UHFFFAOYSA-N 0.000 description 1
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- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 1
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Images
Classifications
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- 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
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
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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
- C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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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
- C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
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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
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
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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
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
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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
- C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of silicon
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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
- C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
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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
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- C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
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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
- C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
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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
- C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
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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
- 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
- C23C8/10—Oxidising
- C23C8/12—Oxidising using elemental oxygen or ozone
- C23C8/14—Oxidising of ferrous surfaces
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- F—MECHANICAL ENGINEERING; LIGHTING; HEATING; WEAPONS; BLASTING
- F16—ENGINEERING ELEMENTS AND UNITS; GENERAL MEASURES FOR PRODUCING AND MAINTAINING EFFECTIVE FUNCTIONING OF MACHINES OR INSTALLATIONS; THERMAL INSULATION IN GENERAL
- F16B—DEVICES FOR FASTENING OR SECURING CONSTRUCTIONAL ELEMENTS OR MACHINE PARTS TOGETHER, e.g. NAILS, BOLTS, CIRCLIPS, CLAMPS, CLIPS OR WEDGES; JOINTS OR JOINTING
- F16B33/00—Features common to bolt and nut
- F16B33/06—Surface treatment of parts furnished with screw-thread, e.g. for preventing seizure or fretting
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- F—MECHANICAL ENGINEERING; LIGHTING; HEATING; WEAPONS; BLASTING
- F16—ENGINEERING ELEMENTS AND UNITS; GENERAL MEASURES FOR PRODUCING AND MAINTAINING EFFECTIVE FUNCTIONING OF MACHINES OR INSTALLATIONS; THERMAL INSULATION IN GENERAL
- F16B—DEVICES FOR FASTENING OR SECURING CONSTRUCTIONAL ELEMENTS OR MACHINE PARTS TOGETHER, e.g. NAILS, BOLTS, CIRCLIPS, CLAMPS, CLIPS OR WEDGES; JOINTS OR JOINTING
- F16B35/00—Screw-bolts; Stay-bolts; Screw-threaded studs; Screws; Set screws
-
- 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
- C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
- C21D1/18—Hardening; Quenching with or without subsequent tempering
-
- 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
- C21D8/005—Modifying the physical properties by deformation combined with, or followed by, heat treatment of ferrous alloys
-
- F—MECHANICAL ENGINEERING; LIGHTING; HEATING; WEAPONS; BLASTING
- F16—ENGINEERING ELEMENTS AND UNITS; GENERAL MEASURES FOR PRODUCING AND MAINTAINING EFFECTIVE FUNCTIONING OF MACHINES OR INSTALLATIONS; THERMAL INSULATION IN GENERAL
- F16B—DEVICES FOR FASTENING OR SECURING CONSTRUCTIONAL ELEMENTS OR MACHINE PARTS TOGETHER, e.g. NAILS, BOLTS, CIRCLIPS, CLAMPS, CLIPS OR WEDGES; JOINTS OR JOINTING
- F16B33/00—Features common to bolt and nut
Definitions
- the present invention relates to a bolt used in automobiles, various industrial machines, etc., a steel for a bolt for realizing the bolt, and a method for producing the bolt, and in particular, excellent hydrogen resistance even at high strength.
- the present invention relates to a high-strength bolt exhibiting embrittlement characteristics, a bolt steel used for manufacturing the bolt, and a method for manufacturing the bolt.
- Patent Documents 1 to 3 show that various alloy elements are adjusted. It is disclosed that a high-strength bolt excellent in delayed fracture characteristics can be obtained.
- Patent Document 4 discloses that many alloy compounds are precipitated by high-temperature tempering after quenching of alloy steel, and traps hydrogen (diffusible hydrogen) moving in the steel in the precipitate. Shows that hydrogen embrittlement resistance can be improved.
- Patent Documents 1 to 4 show excellent hydrogen embrittlement resistance (delayed fracture resistance) in an environment where the amount of hydrogen is relatively small, but hydrogen trapped in carbide is an environment. Is released from trapping sites due to temperature changes in steel and stress fluctuations in steel materials. Therefore, in an environment where there is a large amount of hydrogen where all hydrogen trap sites are consumed or in environments where severe steel corrosion occurs, hydrogen from the trap sites. Is released, the amount of diffusible hydrogen increases, and delayed fracture is more likely to occur.
- JP 60-114551 A JP-A-2-267243 JP-A-3-243745 Japanese Patent No. 4031068
- the present invention has been made in view of the above circumstances, and the purpose thereof is a bolt that exhibits excellent hydrogen embrittlement resistance even at high strength (particularly in an environment with a large amount of hydrogen, Another object of the present invention is to provide a bolt that exhibits excellent hydrogen embrittlement resistance even in an environment involving severe steel corrosion, a bolt steel useful for manufacturing the bolt, and a method for manufacturing the bolt.
- the steel for bolts of the present invention that has solved the above problems is C: 0.30 to 0.50% (meaning “mass%”; the same shall apply hereinafter) Si: 1.0-2.5%, Mn: 0.1 to 1.5% P: 0.015% or less (excluding 0%), S: 0.015% or less (excluding 0%), Cr: 0.15 to 2.4%, Al: 0.010 to 0.10%, and N: 0.001 to 0.10%, Cu: 0.1 to 0.50% and Ni: 0.1 to 1.0%, [Ni] / [Cu] ⁇ 0.5 (where [Ni] indicates the amount of Ni (mass%) in steel) And [Cu] is contained so as to satisfy the Cu content (% by mass) in the steel) Ti: 0.05 to 0.2% and V: 0.2% or less (including 0%), [Ti] + [V]: 0.085 to 0.30% (the above [Ti] is in the steel) Ti amount (mass%) is shown, and the [V] is contained so as to satisfy the V amount (mass%) in
- the present invention has the chemical component,
- the austenite grain size number of the bolt shaft portion is 9.0 or more,
- Bolts having a characteristic in which the G value (%) indicating the proportion of carbide precipitated at the austenite grain boundary of the bolt shaft portion satisfies the following formula (1) are also included.
- G value (L / L0) ⁇ 100 ⁇ 60 (1)
- L the total length of carbides having a thickness of 50 nm or more precipitated at the austenite grain boundaries
- L0 Indicates the length of the austenite grain boundary.
- the bolt preferably has a Fe oxide layer containing Si and Cu on the surface of the bolt shaft portion, and the thickness of the oxide layer is preferably 2.0 to 100 nm.
- the bolt preferably has a tensile strength of 1400 MPa or more.
- the present invention further includes a method for producing a bolt, which uses the steel having the chemical component, performs hot rolling after heating to 1050 ° C. or higher, and sets the finish rolling temperature.
- the tempering is performed at a temperature of 400 ° C. or more and T ° C. or less shown in the following formula (2).
- T (° C.) 68.2 Ln [Si] +480 (2)
- Ln represents the natural logarithm
- [Si] represents the amount of Si in the steel (mass%).
- the tempering is performed using a bolt in which the Fe oxide layer on the surface of the bolt shaft portion is suppressed to 0 to 100 nm, and the atmosphere is inert with an oxygen concentration of 10 ppm (volume basis) or less.
- a gas atmosphere is preferable.
- the strength of austenite grain boundaries which are the starting point of delayed fracture, is increased by appropriately controlling the chemical components of steel and then appropriately adjusting the production conditions, and hydrogen trap sites such as carbides.
- a high-strength bolt that exhibits excellent hydrogen embrittlement resistance is realized not only in an environment with a relatively small amount of hydrogen but also in an environment with a large amount of hydrogen that consumes all hydrogen trap sites. it can.
- the steel for bolts of the present invention is excellent in hot ductility and cold workability (cold forgeability, in particular, bolt forging), the bolt can be manufactured with high productivity.
- FIG. 1 is a graph showing the relationship between the hydrogen embrittlement resistance (results of evaluation of hydrogen embrittlement resistance) and the G value.
- FIG. 2 is a photograph of the austenite grain boundaries of the inventive steel and the comparative steel.
- FIG. 3 is a schematic view showing the shape of a test piece used for measurement of hydrogen embrittlement resistance in a corrosive environment.
- the present inventor is considered to be effective for detoxification of hydrogen, which is one of the factors of delayed fracture phenomenon of high-strength bolts.
- the effect of the hydrogen trap site due to the carbonitride that has been developed was verified anew.
- carbon traps were trapped at these hydrogen trap sites, as described above, although the effect of fixing and detoxifying diffusible hydrogen, which is the main cause of hydrogen embrittlement, is certain. It was found that hydrogen is easily released from the trap site due to changes in the temperature of the steel material and fluctuations in applied stress, and again causes hydrogen embrittlement as diffusible hydrogen.
- the hydrogen released from the hydrogen trap site is fixed again and detoxified if there is another hydrogen trap site around it, but in an environment with a large amount of hydrogen where the hydrogen trap site in the steel becomes saturated. It was confirmed that hydrogen embrittlement easily occurred because the released hydrogen was not fixed again.
- the present inventor has conducted extensive research on a method for improving the hydrogen embrittlement resistance even in an environment with a large amount of hydrogen such that the hydrogen trap sites in the steel are saturated.
- increasing the strength of austenite grain boundaries (hereinafter sometimes simply referred to as “crystal grain boundaries”), which is the starting point of hydrogen embrittlement, is the most effective means.
- crystal grain boundaries As a method for increasing the strength of the austenite grain boundaries, conventionally, a method of dividing carbides precipitated at the grain boundaries by increasing the tempering temperature has been adopted. It has been found that it is effective to prevent precipitation as much as possible, and for that purpose, in quenching and tempering performed after bolt forming, it has been found that performing tempering in a relatively low temperature range is the most effective.
- the tempering temperature when the tempering temperature is less than 400 ° C., the ratio of the yield stress to the tensile strength (yield ratio) decreases, and it becomes difficult to increase the axial force at the time of bolt fastening, and the relaxation characteristics are also improved. There is concern about the decline. Therefore, on the premise that the tempering temperature is 400 ° C. or higher, the composition of steel in which carbides such as cementite hardly precipitate at the austenite grain boundaries even when the tempering temperature is 400 ° C. or higher was examined.
- the carbide precipitation temperature can be shifted to the high temperature side by adding 1.0% or more of Si. This is presumed to be because the presence of Si around the carbon solid-dissolved in the iron after quenching inhibits carbon diffusion during tempering and makes it difficult for carbides to precipitate.
- carbonized_material can be shifted to a high temperature side by containing Si or more than a fixed amount.
- transition carbides such as ⁇ carbide and ⁇ carbide are stabilized, and these transition carbides also have the effect of delaying hydrogen diffusion in steel.
- the apparent hydrogen diffusion coefficient obtained by the hydrogen permeation test is characterized by being as slow as 9.5 ⁇ 10 ⁇ 7 cm 2 / s or less.
- the low hydrogen diffusion coefficient and the slow accumulation of hydrogen at the austenite grain boundaries are also considered to have an effect on improving the hydrogen embrittlement resistance.
- the G value [(L / L0) ⁇ 100] (%) indicating the proportion of carbide precipitated at the austenite grain boundary of the bolt shaft portion is set to satisfy the following formula (1).
- L the total length of carbides having a thickness of 50 nm or more precipitated at the austenite grain boundaries
- L0 Indicates the length of the austenite grain boundary.
- spherical carbides and film-like carbides are not present in the austenite grain boundaries, or the carbides are suppressed to 60% or less with respect to the length of the grain boundaries. Even when the carbide is present on the crystal grain boundary, it is neglected because the adverse effect on the hydrogen embrittlement resistance is low when the thickness of the carbide (length in the direction perpendicular to the crystal grain boundary) is 50 nm or less. it can.
- the G value is preferably 45% or less, more preferably 35% or less. The lower the amount of carbide precipitated at the austenite grain boundaries, the better.
- the lower limit is not particularly limited, but it is usually about 5% or more.
- the bolt of the present invention has an austenite grain size number of the bolt shaft portion of 9.0 or more.
- the austenite grain size number is preferably 9.5 or more, more preferably 10.0 or more.
- the larger the austenite grain size number, the better, and the upper limit is not particularly limited, but it is usually 15 or less.
- the present inventor has only to control the type and thickness of the oxide layer of the bolt surface layer portion generated in the quenching and tempering process during bolt manufacturing. I found. This is because the kind and thickness of the oxide layer strongly affect the hydrogen penetration characteristics into the steel. Specifically, an Fe oxide layer containing Si and Cu and having a thickness of 2.0 to 100 nm is a dense layer excellent in corrosion resistance as an oxide layer of the bolt surface layer portion, It was found that the effect of suppressing hydrogen intrusion was very high.
- the Fe oxide layer of the present invention will be described below.
- the Fe oxide layer of the present invention contains Si and Cu, and mainly contains (Fe, Si) 3 O 4 , (Fe, Cr) 3 O 4, etc. (Note that Cu is mainly Fe. It is thought that it exists alone in the oxide layer).
- the Fe oxide layer of the present invention contains an amount that the spectrum of Si and Cu at an acceleration voltage of 20 kV can be clearly distinguished from noise components when analyzed by EDX (Energy Dispersive X-ray Spectrometer). .
- the Fe oxide layer of the present invention is dense and has a hydrogen penetration inhibiting action.
- the thickness of the oxide layer is preferably 2.0 nm or more. More preferably, it is 5 nm or more, More preferably, it is 10 nm or more.
- the dense Fe oxide layer according to the present invention is very thin, and its upper limit is 100 nm or less.
- C 0.30 to 0.50% C needs to be added to ensure the tensile strength of the steel.
- the C content is determined to be 0.30% or more.
- the amount of C is preferably 0.35% or more, more preferably 0.39% or more.
- the C amount is set to 0.50% or less.
- the amount of C is preferably 0.48% or less, and more preferably 0.45% or less.
- Si acts as a deoxidizer during melting and is an element necessary as a solid solution element for strengthening steel. Further, in the present invention, as described above, an element that suppresses carbides precipitated at the austenite grain boundaries to increase the strength of the grain boundaries, stabilizes transition carbides, and lowers the hydrogen diffusion coefficient in the steel. It is also important. In order to exert such an effect, the Si amount is determined to be 1.0% or more. The amount of Si is preferably 1.2% or more, and more preferably 1.5% or more. On the other hand, when the amount of Si becomes excessive, the cold workability of the steel material is lowered, and the grain boundary oxidation during quenching is promoted to reduce the hydrogen embrittlement resistance. Therefore, the amount of Si is set to 2.5% or less. The amount of Si is preferably 2.3% or less, and more preferably 2.0% or less.
- Mn 0.1 to 1.5% Mn is an element that improves hardenability and is an important element for achieving high strength. Further, since Mn easily forms a compound with S, addition of a certain amount or more also has an effect of suppressing the formation of FeS that precipitates at the grain boundary and causes a decrease in grain boundary strength. In order to effectively exhibit such an action, the amount of Mn is determined to be 0.1% or more. The amount of Mn is preferably 0.13% or more, more preferably 0.15% or more. On the other hand, when the amount of Mn becomes excessive, segregation of MnS to the grain boundary is promoted, the grain boundary strength is lowered, and the hydrogen embrittlement resistance is lowered. Therefore, the amount of Mn is set to 1.5% or less. The amount of Mn is preferably 1.0% or less, more preferably 0.5% or less.
- P 0.015% or less (excluding 0%) P causes grain boundary segregation to lower the grain boundary strength and lower the hydrogen embrittlement resistance. Therefore, the P content is set to 0.015% or less.
- the amount of P is preferably 0.010% or less, and more preferably 0.008% or less. The smaller the amount of P, the better. However, since it causes an increase in the manufacturing cost of the steel material, it is difficult to make it 0%, and a residual of about 0.001% is allowed.
- S 0.015% or less (excluding 0%)
- S is an element that forms sulfide (MnS).
- MnS sulfide
- the S amount is determined to be 0.015% or less.
- the amount of S is preferably 0.010% or less, and more preferably 0.005% or less.
- S is preferably as small as P, but it causes an increase in the manufacturing cost of the steel material. Therefore, it is difficult to set S to 0%, and a residual of about 0.001% is allowed.
- Cr 0.15 to 2.4% Cr is an important element for improving cold forgeability (particularly bolt forgeability) because it becomes a nucleus of spherical carbide formation during spheroidizing annealing and can promote softening. It is also an element that contributes to improving the corrosion resistance of steel in corrosive environments. In order to effectively exhibit such an action, the Cr content is determined to be 0.15% or more. The amount of Cr is preferably 0.5% or more, and more preferably 0.8% or more. On the other hand, when the amount of Cr is excessive, coarse carbonitrides are formed and the toughness is deteriorated. As a result, the hydrogen embrittlement resistance is deteriorated. Therefore, the Cr amount is determined to be 2.4% or less. The amount of Cr is preferably 1.5% or less, more preferably 1.3% or less, and still more preferably 1.2% or less.
- Al 0.010 to 0.10% Al, like Si, functions as a deoxidizer during melting and combines with N in steel to produce AlN, thereby suppressing crystal grain growth, resulting in hydrogen resistance due to refinement of crystal grains. It is an element that can improve the embrittlement characteristics. Therefore, the Al content needs to be 0.010% or more, and more preferably 0.015% or more. On the other hand, when the amount of Al is excessive, oxide inclusions such as Al 2 O 3 are generated, which becomes a stress concentration source and deteriorates the hydrogen embrittlement resistance. In addition, coarse AlN is generated, the crystal grains cannot be refined, the toughness is lowered, and the hydrogen embrittlement resistance is also lowered. Therefore, the Al content is determined to be 0.10% or less. The amount of Al is preferably 0.07% or less, more preferably 0.05% or less.
- N 0.001 to 0.10%
- N is an element that forms nitrides to refine crystal grains and thus improves the resistance to hydrogen embrittlement.
- the N amount needs to be 0.001% or more, more preferably 0.002% or more, and further preferably 0.004% or more.
- the N content is determined to be 0.015% or less.
- the amount of N is preferably 0.007% or less, and more preferably 0.006% or less.
- the amount of Cu is determined to be 0.1% or more.
- the amount of Cu is preferably 0.15% or more, and more preferably 0.20% or more.
- the amount of Cu becomes excessive, the above effects are saturated, and hot ductility is reduced, and the productivity of steel is reduced.
- cold workability, toughness, and hydrogen embrittlement resistance are reduced.
- the hardness of the steel material at the time of bolting increases, resulting in a decrease in the mold life. Therefore, the Cu content is set to 0.50% or less.
- the amount of Cu is preferably 0.4% or less, and more preferably 0.3% or less.
- Ni like Cu, is an element effective for ensuring the strength at low temperature tempering where precipitation of transition carbide occurs. Moreover, there exists an effect
- the amount of Cu and the amount of Ni each satisfy the above ranges, and [Ni] / [Cu] ⁇ 0.5 (where [Ni] indicates the amount of Ni in the steel (mass%), and [Cu] is in the steel)
- the amount of Cu (mass%).
- Ti and V are both elements that have the effect of improving toughness by producing fine carbides and making crystal grains fine.
- the Ti content is determined to be 0.05% or more.
- the amount of Ti is preferably 0.060% or more, and more preferably 0.065% or more.
- V amount becomes like this. Preferably it is 0.1% or more, More preferably, it is 0.14% or more.
- both Ti and V are contained excessively, coarse carbonitrides are formed, and cold forgeability (particularly bolt forgeability) deteriorates.
- the upper limits of Ti and V are each set to 0.2% or less.
- the amount of Ti is preferably 0.15% or less, more preferably 0.1% or less.
- V amount becomes like this. Preferably it is 0.18% or less, More preferably, it is 0.17% or less.
- the total amount of Ti amount and V amount ([Ti] + [V]) is set to 0.085% or more.
- the total amount of Ti and V ((Ti) + [V]) was set to 0.30% or less.
- it is 0.26% or less, More preferably, it is 0.24% or less. Since the grain refinement effect is larger for Ti than V, only Ti is essential in the present invention.
- the basic components of the steel for bolts (bolts) according to the present invention are as described above, and the balance is substantially iron. However, it is naturally allowed that steel contains inevitable impurities brought in depending on the situation of raw materials, materials, manufacturing equipment, and the like.
- the steel for bolts according to the present invention may further include the following Mo as required.
- Mo 0.1% or less (excluding 0%) Mo is an element that improves hardenability and is an effective element for achieving high strength. Further, since it has an effect of suppressing grain boundary oxidation, it is an effective element for steel materials with a large amount of Si added as in the present invention. In order to effectively exhibit such an effect, the Mo content is preferably 0.01% or more, and more preferably 0.03% or more. On the other hand, when the amount of Mo becomes excessive, the steel material cost and the manufacturing cost of the bolt are increased. Therefore, the amount of Mo is preferably 0.1% or less, more preferably 0.08% or less, and still more preferably. 0.07% or less.
- the bolt according to the present invention is prepared by melting steel having the above chemical components in accordance with a normal melting method, casting, hot rolling, wire drawing, softening treatment such as spheroidizing annealing, descaling and finish wire drawing. Thereafter, it can be manufactured by bolting by cold forging or cold forging, and further by quenching and tempering.
- a steel material having the above chemical composition By using a steel material having the above chemical composition, a bolt with high strength and excellent hydrogen embrittlement resistance can be obtained, but in order to obtain a bolt with high strength and further excellent hydrogen embrittlement resistance. It is important to manufacture by appropriately controlling the conditions of hot rolling and quenching and tempering treatment in the series of steps described above. It is also preferable to control the quenching conditions.
- the steel for bolts and the method for producing the bolts will be described.
- the pre-structure is generally a spheroidized structure, but in order to make the carbide dispersion uniform in the spheroidized structure, the heating temperature before hot rolling (reheating temperature in billet) and the hot rolling conditions are set.
- the heating temperature before hot rolling reheating temperature in billet
- the hot rolling conditions are set.
- carbides such as Cr, Ti, and V, which are the cores of spheroidized carbides, and the number of triple points (points where three austenite crystal grains are in contact) of austenite grain boundaries It is important to increase, that is, to refine the austenite crystal grains.
- the heating temperature before hot rolling should be 1050 ° C. or higher, and Cr, Ti, V, etc. must be dissolved in the austenite region.
- the heating temperature is preferably 1100 ° C. or higher, more preferably 1150 ° C. or higher.
- the upper limit temperature is about 1300 ° C. from the viewpoint of manufacturing cost.
- the finish rolling temperature is preferably 950 ° C. or lower, and more preferably 900 ° C. or lower.
- the lower limit is preferably set to 700 ° C.
- the finish rolling temperature is the average surface temperature that can be measured with a radiation thermometer before the final rolling pass or before the rolling roll group.
- the heating temperature during quenching is preferably 860 to 930 ° C. If the heating temperature at the time of quenching is too low, the carbide generated by hot rolling or softening treatment does not sufficiently dissolve, so the strength decreases, and if coarse carbides remain at the austenite grain boundaries, hydrogen embrittlement occurs. Therefore, the hydrogen embrittlement resistance deteriorates.
- a more preferable heating temperature is 880 ° C. or higher, and further preferably 890 ° C. or higher.
- the quenching temperature is too high, the crystal grains are coarsened and the hydrogen embrittlement resistance is deteriorated. More preferably, it is 920 degrees C or less, More preferably, it is 910 degrees C or less.
- the atmosphere at the time of quenching is not particularly limited, but from the viewpoint of manufacturing cost, it is desirable to treat in a normal atmospheric atmosphere.
- the corrosion resistance is improved by forming a dense Fe oxide layer on the surface of the bolt shaft portion, there is no Fe oxide layer on the bolt shaft portion before tempering, or Fe Even if an oxide layer is present, it is important that the oxide layer be suppressed to 100 nm or less. This is because if a sparse Fe oxide layer exceeding 100 nm exists before tempering, formation of the dense oxide layer of the present invention is prevented during tempering.
- the atmosphere during quenching may be a condition that does not form a sparse Fe oxide layer, that is, an inert gas atmosphere with a reduced oxygen concentration.
- a sparse oxide layer is formed.
- the formed sparse oxide layer may be removed.
- the method for removing the sparse oxide layer is not particularly limited, and may be pickled or removed mechanically, for example.
- a heating method You may implement by a normal electric furnace and gas furnace, and you may implement by high frequency heating.
- the tempering temperature is 400 ° C. or higher and T ° C. or lower shown in the following formula (2).
- T (° C.) 68.2 Ln [Si] +480 (2)
- Ln represents the natural logarithm
- [Si] represents the amount of Si in the steel (mass%).
- the precipitation temperature of carbide during tempering varies depending on the amount of Si in the steel, and is represented by (68.2 Ln [Si] +480) ° C. If it is higher than this temperature, carbide precipitates at the crystal grain boundaries, and the hydrogen embrittlement resistance deteriorates due to a decrease in grain boundary strength.
- tempering is performed at a temperature of (68.2Ln [Si] +480) ° C. or lower.
- it is (T-20) ° C. or lower, that is, (68.2 Ln [Si] +460) ° C. or lower, more preferably (T-40) ° C. or lower, that is, (68.2 Ln [Si] +440) ° C. or lower.
- tempering temperature is too low, the yield ratio decreases as described above, and it is difficult to fasten the bolt with a high axial force. Therefore, tempering is performed at 400 ° C. or higher.
- it is 420 degreeC or more, More preferably, it is 425 degreeC or more.
- the atmosphere at the time of tempering is not particularly limited, but from the viewpoint of production cost, it is desirable to perform the treatment in a normal air atmosphere.
- the atmosphere is preferably an inert gas atmosphere having an oxygen concentration of 10 ppm (volume basis) or less.
- the inert gas atmosphere having an oxygen concentration of 10 ppm (volume basis) or less.
- N 2 or argon can be used as the inert gas. It does not specifically limit regarding a heating method, You may implement by a normal electric furnace and gas furnace, and you may implement by high frequency heating.
- Cooling condition Oil cooling or water cooling [Tempering condition] Holding time after heating: 10 minutes or more (more preferably 20 minutes or more), 90 minutes or less (more preferably 45 minutes or less) Cooling condition: oil cooling or water cooling
- the bolt of the present invention has particularly high tensile strength of 1400 MPa or more, and exhibits excellent hydrogen embrittlement resistance even at such high strength.
- Example 1 In Example 1, the result of investigating the hydrogen embrittlement resistance in the cathode charge environment will be described.
- a steel having the chemical components shown in Table 1 (the balance being iron and inevitable impurities) was melted and cast according to a normal melting method, and then hot-rolled under the conditions shown in Tables 2 and 3 to obtain a rolled material having a diameter of 14 mm Got.
- the rolled material was subjected to wire drawing and spheroidizing annealing after descaling / coating treatment, and finishing wire drawing after further descaling / coating treatment.
- a flange bolt of M12 x 1.25P and length 200mmL was made by cold forging using a parts former (manufactured by Sakamura Machinery Co., Ltd .: NBP550).
- quenching and tempering were performed under the conditions shown in Table 2 or Table 3 below.
- quenching heating time 30 minutes
- quenching furnace atmosphere air
- quenching cooling condition oil cooling (70 ° C.)
- tempering heating time 45 minutes
- tempering furnace atmosphere Air and tempering cooling conditions: Oil cooling (25 ° C.).
- the hydrogen diffusion coefficient was measured by cutting the shaft part of the bolt in a cross section, performing mechanical polishing and then electrolytic polishing, processing it into a 0.1 mm-thick sheet material, and then performing electrochemical Measured by a static hydrogen permeation test. Specifically, both surfaces of the plate were used as hydrogen generation / hydrogen abstraction surfaces, and the rate (time-dependent change) of hydrogen atoms generated by electrolysis of the acid solution through the plate was evaluated. The permeated hydrogen flux was detected by measuring the oxidation current in the hydrogen drawn surface / alkaline solution.
- Test solution Hydrogen detection side 1N-NaOH Hydrogen generation side 0.5 mol / L H 2 SO 4 + 0.01 mol / L KSCN Electrolysis conditions: Hydrogen detection side 150 mV vs Ag / AgCl Hydrogen generation side BuildUp 5 mA / cm 2 , Decay 1 mA / cm 2 Test area: 0.18 cm 2
- Tables 1 to 3 can be considered as follows (the following numbers indicate “Experiment No.” in Tables 2 and 3). That is, no. In Nos. 1 to 12, since the component composition of steel and the production conditions are appropriately controlled, all of them achieve high strength of 1400 MPa or more and realize excellent hydrogen embrittlement resistance.
- No. Nos. 13 to 34 are inferior in strength or hydrogen embrittlement resistance because at least one of the chemical component composition and production conditions of the steel was inappropriate.
- No. No. 13 is an example having a low reheating temperature. Since coarse carbonitrides such as Cr, Ti, and V remained in the steel, the subsequent spheroidizing annealing became insufficient and the bolt forging decreased.
- No. No. 14 is an example in which the tempering temperature is lower than 400 ° C., and since the yield ratio was low, the characteristics as a bolt could not be satisfied.
- No. Nos. 15 to 18 are examples in which the tempering temperature is higher than the T value, and a large amount of carbides precipitated on the crystal grain boundaries, so that the grain boundary strength was lowered and the hydrogen embrittlement resistance was lowered.
- No. Nos. 19 to 21 are examples in which it is difficult to simultaneously satisfy the tensile strength, the yield ratio, and the hydrogen embrittlement resistance because a plurality of additive elements do not satisfy the requirements of the present invention.
- No. 19 is an example using SCM435, which is a JIS standard steel, but since the amount of Si is insufficient, the T value is lower than 400 ° C., and the yield ratio and hydrogen embrittlement resistance could not be satisfied at the same time. .
- tempering was performed at 480 ° C., but the tensile strength was below 1400 MPa.
- No. 22 is an example with little C, and even when tempered at 400 ° C., the tensile strength was less than 1400 MPa.
- No. No. 23 is an example with little Si, and the T value was lower than 400 ° C., and the yield ratio and the resistance to hydrogen embrittlement could not be satisfied at the same time.
- tempering was performed at 400 ° C., but a large number of carbides precipitated at the austenite grain boundaries, so that the hydrogen embrittlement resistance could not be satisfied.
- No. No. 24 is an example with less Mn. 25 is an example with much Mn.
- FeS was generated in a part of the crystal grain boundary, and the grain boundary strength was lowered, so that the hydrogen embrittlement resistance was lowered.
- No. 25 the grain boundary segregation of MnS resulted in a decrease in grain boundary strength, and the hydrogen embrittlement resistance decreased.
- No. No. 26 is an example with less Cu. 27 is an example where there is much Cu. No. In No. 26, even after tempering at 400 ° C., the tensile strength was less than 1400 MPa, so the tensile strength and the yield ratio could not be satisfied simultaneously. No. In No. 27, bolt forging decreased due to a decrease in toughness.
- No. No. 28 is an example in which the value of Ni / Cu is small, and it is considered that the resistance to hydrogen embrittlement deteriorated due to the occurrence of hot embrittlement during continuous casting and the presence of microcracks and the like inside.
- No. No. 29 is an example with little Cr.
- 30 is an example with much Cr.
- No. No. 31 is an example with a large amount of V. Coarse V carbonitride was formed, and bolt forging was deteriorated, so that bolt forging could not be performed.
- No. 32 is an example in which the total amount of V and Ti is large. A large amount of hydrogen trap sites are generated by tempering, a large amount of hydrogen is fixed to the trap sites, and hydrogen is released when the temperature changes during the test. The hydrogen embrittlement resistance is considered to have deteriorated.
- No. No. 33 is an example of a large amount of Al, and since coarse AlN was produced, the crystal grains became coarse during quenching, the toughness was lowered, and the hydrogen embrittlement resistance was lowered.
- No. No. 34 is an example with a large amount of N. Since the amount of N dissolved in the steel increased, the bolt forging deteriorated.
- FIG. 1 is a graph showing the relationship between the hydrogen embrittlement resistance (evaluation result of hydrogen embrittlement resistance) and the G value using the results of Example 1. From FIG. 1, there is a correlation between the hydrogen embrittlement resistance (results of evaluation of hydrogen embrittlement resistance) and the G value, and the hydrogen embrittlement resistance improves as the G value decreases. It can be seen that the G value should be 60 or less in order to make the embrittlement value 0.60 or more.
- FIG. 2 shows experiment no. 1 and Experiment No. It is the photograph which each image
- Example 2 In Example 2, the results of investigating hydrogen embrittlement resistance in a corrosive environment will be described.
- Example 2 After the steel having the chemical components shown in Table 4 (the balance being iron and inevitable impurities) was melted and cast according to a normal melting method, hot rolling (in Example 2, the reheating temperature in each example: 1050 ° C. or higher, finish rolling temperature: 1000 ° C. or lower) to obtain a rolled material having a diameter of 14 mm. Delayed fracture specimens and tensile specimens were collected from the rolled material by cutting and quenched and tempered under the conditions shown in Tables 5 and 6. Regarding other quenching and tempering conditions, quenching heating time: 30 minutes, quenching furnace atmosphere: N 2 , quenching cooling condition: oil cooling (70 ° C.), tempering time: 1 hour, tempering cooling condition: air cooling did.
- Component elements of the furnace atmosphere in the quenching and tempering was set to N 2, to control the atmosphere by the following procedure.
- N 2 gas 1.013 ⁇ 10 5 Pa.
- the oxygen concentration in the atmosphere is about 21% by volume
- the oxygen concentration is calculated as 0.8 ppm (volume basis).
- Example 2 (4) Evaluation of cold forgeability
- the following cold compression process was performed simulating bolting and the cold forgeability was evaluated. That is, the rolled material was subjected to spheroidizing annealing (held at 780 ° C. for 6 hours, cooling rate 10 ° C./hour) to produce a cold compression test piece ( ⁇ 10 mm ⁇ L15 mm), and whether or not cracking occurred during 70% compression Cold forgeability was evaluated.
- the strain rate during cold compression was 10 s ⁇ 1 .
- the weight change amount of the test piece before and after acid immersion was measured, and the value obtained by dividing the weight change amount by the weight of the test piece before acid immersion was multiplied by 100 to obtain the corrosion weight loss (%).
- those having a delayed fracture strength ratio of 0.70 or more were evaluated as having excellent hydrogen embrittlement resistance in a corrosive environment.
- Tables 4 to 6 can be considered as follows (the following No. indicates “Experiment No.” in Tables 5 and 6). That is, no. For 1-3, 6-10, the steel component composition and manufacturing conditions are appropriately adjusted, so that both achieve a high strength of 1400 MPa or higher and a delayed fracture strength ratio of 0.70 or higher, corrosive environment It also has excellent hydrogen embrittlement resistance.
- No. Nos. 4 to 5 and 11 to 26 resulted in inferior hydrogen embrittlement resistance in a strength or corrosive environment because at least one of the chemical component composition and production conditions of the steel was inappropriate.
- No. 5 shows the effect of the heating temperature during quenching. No. In No. 5, since the heating temperature at the time of quenching was high, the crystal grains were coarsened and the toughness was lowered, so that the hydrogen embrittlement resistance in a corrosive environment was lowered.
- No. No. 17 is an example in which the amount of Cu is large, and the cold forgeability deteriorated due to the decrease in toughness.
- No. No. 18 is an example in which the value of Ni / Cu is small, and it is considered that the resistance to hydrogen embrittlement in a corrosive environment has decreased due to the decrease in hot ductility and the presence of microcracks and the like inside.
- No. No. 19 is an example with a small amount of Si, and the desired tensile strength could not be ensured at a tempering temperature of 425 ° C.
- No. No. 20 is an example of a large amount of C.
- the delayed fracture susceptibility increases, so the hydrogen embrittlement resistance in a corrosive environment has decreased.
- the cold forgeability also fell.
- No. No. 21 is an example with a small amount of Mn. 22 is an example with a large amount of Mn.
- FeS was generated in a part of the crystal grain boundary, and the grain boundary strength was lowered, so that the hydrogen embrittlement resistance in a corrosive environment was lowered. Moreover, the cold forgeability also fell. Furthermore, no. In No. 22, grain boundary segregation of MnS resulted in a decrease in grain boundary strength, resulting in a decrease in hydrogen embrittlement resistance in a corrosive environment.
- No. No. 23 is an example in which the amount of Cr is small. Since the thickness of the Fe oxide layer is insufficient and the corrosion resistance is lowered, the hydrogen embrittlement resistance in a corrosive environment is lowered.
- No. No. 24 is an example with a large amount of Cr, and coarse carbonitrides formed after tempering, and the toughness was lowered, so the hydrogen embrittlement resistance in a corrosive environment was lowered.
- No. No. 25 is an example in which the amount of Al is large and the quenching temperature is too high. Since coarse AlN was generated, it became a stress concentration source and deteriorated the hydrogen embrittlement resistance in a corrosive environment.
- No. No. 26 is an example in which the amount of N is large. Since the amount of N dissolved in steel increased, the hydrogen embrittlement resistance in a corrosive environment deteriorated.
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Abstract
Description
C:0.30~0.50%(「質量%」の意味。以下同じ)、
Si:1.0~2.5%、
Mn:0.1~1.5%、
P:0.015%以下(0%を含まない)、
S:0.015%以下(0%を含まない)、
Cr:0.15~2.4%、
Al:0.010~0.10%、および
N:0.001~0.10%を含有し、
Cu:0.1~0.50%およびNi:0.1~1.0%を、[Ni]/[Cu]≧0.5(前記[Ni]は鋼中Ni量(質量%)を示し、前記[Cu]は鋼中Cu量(質量%)を示す)を満たすように含有するとともに、
Ti:0.05~0.2%およびV:0.2%以下(0%を含む)を、[Ti]+[V]:0.085~0.30%(前記[Ti]は鋼中Ti量(質量%)を示し、前記[V]は鋼中V量(質量%)を示す)を満たすように含有し、
残部が鉄および不可避的不純物であるところに特徴を有する。
前記ボルト用鋼は、更に、Mo:0.1%以下(0%を含まない)を含んでいてもよい。
ボルト軸部のオーステナイト結晶粒度番号が9.0以上であり、
ボルト軸部のオーステナイト結晶粒界に析出した炭化物の割合を示すG値(%)が、下記式(1)を満たすところに特徴を有するボルトも含まれる。
G値:(L/L0)×100≦60 …(1)
(式(1)において、
L:オーステナイト結晶粒界に析出した厚さ50nm以上の炭化物の合計長さ、
L0:オーステナイト結晶粒界の長さを示す。)
T(℃)=68.2Ln[Si]+480 …(2)
(式(2)において、Lnは自然対数を示し、[Si]は鋼中Si量(質量%)を示す。)
(L/L0)×100≦60 …(1)
(式(1)において、
L:オーステナイト結晶粒界に析出した厚さ50nm以上の炭化物の合計長さ、
L0:オーステナイト結晶粒界の長さを示す。)
Cは、鋼の引張強度を確保するために添加する必要がある。高強度(特には、引張強度1400MPa以上)を確保するため、C量は0.30%以上と定めた。C量は、好ましくは0.35%以上であり、より好ましくは0.39%以上である。一方、C量が過剰になると、靭性の低下を招くと共に、オーステナイト結晶粒界に炭化物が生成し易くなり粒界強度の低下が生じて、耐水素脆化特性が劣化する。更には、冷間加工性(冷間鍛造性、特にはボルト圧造性)の低下も生じる。また、腐食環境ではC量が過剰になると耐食性が悪化する。そこでC量を0.50%以下と定めた。C量は、好ましくは0.48%以下であり、より好ましくは0.45%以下である。
Siは、溶製時の脱酸剤として作用するとともに、鋼を強化する固溶元素として必要な元素である。また本発明においては、上述の通り、オーステナイト結晶粒界に析出する炭化物を抑制して該結晶粒界の強度を高くするとともに、遷移炭化物を安定化させ、鋼中の水素拡散係数を低下させる元素としても重要である。このような作用を発揮させるため、Si量は1.0%以上と定めた。Si量は好ましくは1.2%以上であり、より好ましくは1.5%以上である。一方、Si量が過剰になると、鋼材の冷間加工性が低下するとともに、焼入れ時における粒界酸化を助長して耐水素脆化特性を低下させる。そこで、Si量は2.5%以下と定めた。Si量は、好ましくは2.3%以下であり、より好ましくは2.0%以下である。
Mnは、焼入れ性向上元素であり、高強度化を達成する上で重要な元素である。また、MnはSと化合物を形成しやすいため、一定以上添加することにより、結晶粒界に析出して粒界強度の低下を招くFeSの生成を抑制する効果も有する。このような作用を有効に発揮させるため、Mn量は0.1%以上と定めた。Mn量は、好ましくは0.13%以上であり、より好ましくは0.15%以上である。一方、Mn量が過剰になると、粒界へのMnSの偏析を助長して粒界強度が低下し耐水素脆化特性が低下する。そこで、Mn量を1.5%以下と定めた。Mn量は、好ましくは1.0%以下であり、より好ましくは0.5%以下である。
Pは、粒界偏析を起こして粒界強度を低下させ、耐水素脆化特性を低下させる。そこで、P量は0.015%以下と定めた。P量は好ましくは0.010%以下であり、より好ましくは0.008%以下である。P量は少なければ少ないほど好ましいが、鋼材の製造コストの増加を招くため、0%とすることは難しく、0.001%程度の残存は許容される。
Sは、硫化物(MnS)を形成する元素である。S量が過剰になると、前記MnSとして粗大なものが形成され、この粗大なMnSが応力集中箇所となって耐水素脆化特性の低下を招く。そこでS量は、0.015%以下と定めた。S量は、好ましくは0.010%以下であり、より好ましくは0.005%以下である。Sは、Pと同様に少なければ少ないほど好ましいが、鋼材の製造コストの増加を招くため、0%とすることは難しく、0.001%程度の残存は許容される。
Crは、球状化焼鈍時に球状炭化物形成の核となり、軟化を促進させることができるため、冷間鍛造性(特には、ボルト圧造性)を向上する上で重要な元素である。また、腐食環境での鋼の耐食性向上にも寄与する元素である。このような作用を有効に発揮させるため、Cr量は0.15%以上と定めた。Cr量は、好ましくは0.5%以上であり、より好ましくは0.8%以上である。一方、Cr量が過剰になると、粗大な炭窒化物が形成して靭性が劣化し、その結果、耐水素脆化特性が劣化する。そこで、Cr量を2.4%以下と定めた。Cr量は、好ましくは1.5%以下であり、より好ましくは1.3%以下、更に好ましくは1.2%以下である。
Alは、Siと同様に溶製時の脱酸剤として機能するとともに、鋼中のNと結合しAlNを生成することによって、結晶粒成長を抑制し、結果として結晶粒の微細化により耐水素脆化特性を向上させることのできる元素である。よって、Al量は0.010%以上とする必要があり、より好ましくは0.015%以上である。一方、Al量が過剰になると、Al2O3などの酸化物系介在物を生成し、応力集中源となって耐水素脆化特性を低下させる。また、粗大なAlNが生成して、結晶粒の微細化が図れず、靭性が低下して耐水素脆化特性の低下も生じる。そこで、Al量は0.10%以下と定めた。Al量は、好ましくは0.07%以下であり、より好ましくは0.05%以下である。
Nは、窒化物を形成して結晶粒を微細化し、ひいては耐水素脆化特性を向上させる元素である。このような作用を有効に発揮させるため、N量は0.001%以上とする必要があり、より好ましくは0.002%以上であり、さらに好ましくは0.004%以上である。一方、N量が過剰になると、鋼中に固溶するN量が増大し、冷間加工性および耐水素脆化特性を低下させる。従って、N量は0.015%以下と定めた。N量は、好ましくは0.007%以下であり、より好ましくは0.006%以下である。
Cuは、遷移炭化物の析出が起きる低温焼戻しでの強度を確保するのに有効な元素である。また、腐食環境での鋼の耐食性を向上することもできる。このような作用を有効に発揮させるため、Cu量は0.1%以上と定めた。Cu量は、好ましくは0.15%以上であり、より好ましくは0.20%以上である。一方、Cu量が過剰になると、前記効果が飽和するとともに、熱間延性が低下して鋼の生産性が低下する。また冷間加工性の低下や靭性の低下、耐水素脆化特性の低下を招く。更に、ボルト加工時の鋼材硬さが増加して金型寿命の低下ももたらす。そこでCu量は0.50%以下と定めた。Cu量は、好ましくは0.4%以下であり、より好ましくは0.3%以下である。
TiおよびVはいずれも、微細な炭化物を生成し、結晶粒を微細化することで靭性を向上させる効果を有する元素である。このような効果を有効に発揮させるため、Ti量は0.05%以上と定めた。Ti量は、好ましくは0.060%以上であり、より好ましくは0.065%以上である。またV量は、好ましくは0.1%以上であり、より好ましくは0.14%以上である。一方、TiおよびVはいずれも、過剰に含まれると粗大な炭窒化物を形成し、冷間鍛造性(特にはボルト圧造性)が劣化する。また、TiおよびVはいずれも、過剰に含まれると水素トラップサイトが増加して鋼中の水素量が増加し、温度変化や応力変動等によりトラップサイトから水素が開放された際、水素脆化を起こしやすくなる。よって本発明では、TiおよびVの上限をそれぞれ0.2%以下と定めた。Ti量は、好ましくは0.15%以下であり、より好ましくは0.1%以下である。またV量は、好ましくは0.18%以下であり、より好ましくは0.17%以下である。さらにTiとVの結晶粒微細化の効果を有効に発揮させるため、本発明ではTi量とV量の合計量([Ti]+[V])を0.085%以上と定めた。好ましくは0.1%以上であり、より好ましくは0.2%以上である。一方、TiとVによる水素トラップサイトの効果を低減するため、前記Ti量とV量の合計量([Ti]+[V])は0.30%以下とした。好ましくは0.26%以下であり、より好ましくは0.24%以下である。尚、結晶粒微細化効果はVよりもTiの方が大きいため、本発明においてはTiのみを必須としている
Moは、焼入れ性向上元素であり、高強度を達成するのに有効な元素である。また、粒界酸化抑制効果を有しているため、本発明のようにSiの添加量が多い鋼材には有効な元素である。このような効果を有効に発揮させるため、Mo量は0.01%以上含有させることが好ましく、より好ましくは0.03%以上である。一方、Mo量が過剰になると、鋼材コストおよびボルトの製造コストの増加をもたらすため、Mo量は0.1%以下とするのが好ましく、より好ましくは0.08%以下であり、さらに好ましくは0.07%以下である。
T(℃)=68.2Ln[Si]+480 …(2)
(式(2)において、Lnは自然対数を示し、[Si]は鋼中Si量(質量%)を示す。)
焼戻しでの炭化物の析出温度は、鋼中Si量によって変化し、(68.2Ln[Si]+480)℃で表される。この温度より高くなると、結晶粒界に炭化物が析出し、粒界強度の低下により耐水素脆化特性が劣化する。よって焼戻しは、(68.2Ln[Si]+480)℃以下の温度で行う。好ましくは(T-20)℃以下、即ち(68.2Ln[Si]+460)℃以下であり、より好ましくは(T-40)℃以下、即ち(68.2Ln[Si]+440)℃以下である。一方、焼戻し温度を低くしすぎると、前述した通り降伏比が低下し、ボルトを高い軸力で締結することが困難となる。そのため400℃以上で焼戻し処理を行う。好ましくは420℃以上であり、より好ましくは425℃以上である。
加熱後の保持時間:5分以上(より好ましくは15分以上)、60分以下(より好ましくは30分以下)
冷却条件:油冷または水冷
〔焼戻し条件〕
加熱後の保持時間:10分以上(より好ましくは20分以上)、90分以下(より好ましくは45分以下)
冷却条件:油冷または水冷
実施例1では、陰極チャージ環境での耐水素脆化特性を調査した結果を説明する。
ボルトの軸部を横断面(ボルトの軸に対して垂直な断面。以下同じ)で切断後、D/4位置(Dは軸部の直径)の任意の0.039mm2の領域を光学顕微鏡で観察し(倍率:400倍)、JIS G0551に従って結晶粒度番号を測定した。測定は4視野について行い、これらの平均値をオーステナイト結晶粒度とした。
ボルトの引張強度はJIS B1051に従って引張試験を行い求めた。また、降伏比は0.2%耐力を引張強度で除すことで求めた。
上記引張試験で得られた引張強度が1400MPa以上でかつ降伏比が0.90以上のものを対象に、オーステナイト結晶粒界に析出した炭化物の観察を行った。尚、表3におけるNo.15とNo.19は、引張強度が1400MPaに満たない例であるが、参考までに、このオーステナイト結晶粒界に析出した炭化物の観察も行った。
オーステナイト結晶粒界に析出した炭化物は、上記ボルトの軸部を横断面で切断後、集束イオンビーム加工装置(FIB:Focused Ion Beam Process、日立製作所製:FB-2000A)により薄膜試験片を作成し、透過型電子顕微鏡(日立製作所製、JEMS-2100F)を用いて1試料につき3枚ずつ、倍率15万倍でオーステナイト結晶粒界を撮影し、画像解析で結晶粒界に析出した炭化物の長さと厚さ(厚さは、オーステナイト結晶粒界に対して垂直方向の長さ)を算出した。そして、オーステナイト結晶粒界に析出した厚さ50nm以上の炭化物の長さ(L)をオーステナイト結晶粒界の長さ(L0)で除し、百分率で表すことにより、オーステナイト粒界上の炭化物の占有率(G値)を求め、3枚の写真の平均値を表2および表3に記載した。但し、L0は、取得した画像に含まれているすべてのオーステナイト結晶粒界の長さの総和であり、Lは取得した画像に含まれているすべての「オーステナイト結晶粒界に析出した厚さ50nm以上の炭化物」の長さの総和である
水素拡散係数は、上記ボルトの軸部を横断面で切断後、機械研磨、次いで電解研磨を実施し、0.1mm厚さの薄板材に加工した後、電気化学的水素透過試験にて測定した。具体的には、板材の両表面を水素発生/水素引抜面とし、酸溶液の電気分解により生成した水素原子が板材を透過する速度(経時変化)を評価した。透過水素流束の検出は、水素引抜面・アルカリ溶液中での酸化電流の測定によった。試験は室温条件で実施し、水素発生~定常透過(Build Up)、水素発生停止~水素透過停止(Decay)の各過程カーブを3回取得し、拡散係数は拡散方程式の理論解に対して測定結果をカーブフィッティングすることにより実施した。得られたBuild Up、Decayでの各3回、計6回の測定結果を平均し、水素拡散係数とした。なお、実験に用いた溶液、電解条件、試験面積を以下に示す。
試験液:水素検出側 1N-NaOH
水素発生側 0.5mol/L H2SO4 + 0.01mol/L KSCN
電解条件:水素検出側 150mV vs Ag/AgCl
水素発生側 BuildUp 5mA/cm2、Decay 1mA/cm2
試験面積:0.18cm2
耐水素脆化特性の評価は、上記ボルトの軸部に応力集中係数が3となる切欠きを機械加工によって作成し、切欠き底の面積に対して1500MPaの定荷重を負荷後、ただちに切欠き部に陰極チャージを施した。また、最初の3時間は室温(25℃)で実施したが、次の3時間はヒーターで溶液部を加熱することにより50℃で試験を実施し、その後は室温と50℃を3時間毎に繰返し実施した。結果は破断までの時間を測定することにより評価し、破断が生じなかった場合は120時間で試験を打ち切った。試験結果は破断までの時間を120で除した値を耐水素脆化値とし、この耐水素脆化値が0.60以上の場合を耐水素脆化特性に優れているとした。尚、実験に用いた溶液、電解条件は以下の通りである。
試験液:pH3-H2SO4 + 0.01mol/L KSCN
電解電流密度:0.01mA/cm2
実施例2では、腐食環境での耐水素脆化特性を調査した結果を説明する。
引張試験片(JIS14A号)を試験片の長手方向に垂直な断面(横断面)で切断後、D/4位置(Dは軸部の直径)の任意の0.039mm2の領域を光学顕微鏡で観察し(倍率:400倍)、JIS G0551に従って結晶粒度番号を測定した。測定は4視野について行い、これらの平均値をオーステナイト結晶粒度とした。
引張試験は、引張試験片(JIS14A号)を用い、上記JIS Z2241に従って引張強度を測定した。また、前記したオーステナイト結晶粒度測定と同じ領域の硬度を、ビッカース硬度計で測定(荷重:10kg)した。測定は4箇所で行い、これらの平均値をボルト軸部の硬さとした。
軸部の表面酸化層の分析は、上記試験片を横断面(軸心に垂直な断面)で切断して樹脂に埋め込み、まず、電界放射型走査電子顕微鏡(日立製作所製、S-4500)を用いて倍率500倍で表面全周を観察し、熱処理後の軸部に特異箇所のないことと、100nmを超える酸化層の生成がないことを確認した。その後、透過型電子顕微鏡(日立製作所製、JEMS-2100F)で倍率30万倍と150万倍で確認した90°毎に2箇所を写真撮影し(150万倍)、画像解析で表面酸化層の面積を算出した。算出した酸化層の面積を、酸化層直下の地鉄層の長さで除すことによって酸化層の平均厚さとし、2箇所の平均値を求めた。更に酸化物層の組成について、EDX分析を行い、酸化物層に含有される元素の分析を行った。表5および表6において、「Si、Cu含有」が「○」と示されているものは、加速電圧20kVでのSi、Cuのスペクトルが、ノイズ成分と明らかに区別できる量を含有していた。
本実施例2では、ボルト加工を模擬して下記の冷間圧縮加工を行い、冷間鍛造性を評価した。即ち、前記圧延材を球状化焼鈍(780℃で6時間保持、冷却速度10℃/時間)し、冷間圧縮試験片(φ10mm×L15mm)を作製し、70%圧縮時の割れ発生の有無によって冷間鍛造性を評価した。なお、冷間圧縮時のひずみ速度は10s-1とした。
腐食環境での耐水素脆化特性の測定に用いる試験片は、ねじ部の応力集中を模擬できるように、図3に示すような切欠きを設けたものとした。前記試験片を15%HCl溶液に30分浸漬し、水洗・乾燥させた後、一定荷重を付加し、100時間以上破断しない最大荷重を測定した。そして、酸浸漬後に100時間以上破断しない最大荷重を、酸浸漬前に引張試験した際の破断荷重で除した値を、遅れ破壊強度比とし、腐食環境での耐水素脆化特性を評価した。また、酸浸漬前後の試験片の重量変化量を測定し、重量の変化量を酸浸漬前の試験片の重量で除した値に100を掛けたものを腐食減量(%)とした。そして特に、前記遅れ破壊強度比が0.70以上のものを、腐食環境での耐水素脆化特性に優れていると評価した。
Claims (7)
- C:0.30~0.50%(「質量%」の意味。以下同じ)、
Si:1.0~2.5%、
Mn:0.1~1.5%、
P:0.015%以下(0%を含まない)、
S:0.015%以下(0%を含まない)、
Cr:0.15~2.4%、
Al:0.010~0.10%、および
N:0.001~0.10%を含有し、
Cu:0.1~0.50%およびNi:0.1~1.0%を、[Ni]/[Cu]≧0.5(前記[Ni]は鋼中Ni量(質量%)を示し、前記[Cu]は鋼中Cu量(質量%)を示す)を満たすように含有するとともに、
Ti:0.05~0.2%およびV:0.2%以下(0%を含む)を、[Ti]+[V]:0.085~0.30%(前記[Ti]は鋼中Ti量(質量%)を示し、前記[V]は鋼中V量(質量%)を示す)を満たすように含有し、
残部が鉄および不可避的不純物であることを特徴とするボルト用鋼。 - 更に、Mo:0.1%以下(0%を含まない)を含有する請求項1に記載のボルト用鋼。
- 請求項1に記載の化学成分を有し、
ボルト軸部のオーステナイト結晶粒度番号が9.0以上であり、
ボルト軸部のオーステナイト結晶粒界に析出した炭化物の割合を示すG値(%)が、下記式(1)を満たすことを特徴とするボルト。
G値:(L/L0)×100≦60 …(1)
(式(1)において、
L:オーステナイト結晶粒界に析出した厚さ50nm以上の炭化物の合計長さ、
L0:オーステナイト結晶粒界の長さを示す。) - 前記ボルト軸部の表面にSiおよびCuを含有するFe酸化層を有し、該酸化層の厚さが2.0~100nmである請求項3に記載のボルト。
- 引張強度が1400MPa以上である請求項3に記載のボルト。
- 請求項1に記載の化学成分を有する鋼を用い、
熱間圧延を、1050℃以上に加熱してから行い、かつ仕上げ圧延温度を1000℃以下とし、更に、
ボルト成形後に行う焼入れ焼戻しにおいて、焼戻しを、400℃以上であって下記式(2)に示すT℃以下の温度で行うことを特徴とするボルトの製造方法。
T(℃)=68.2Ln[Si]+480 …(2)
(式(2)において、Lnは自然対数を示し、[Si]は鋼中Si量(質量%)を示す。) - 前記焼戻しは、ボルト軸部の表面のFe酸化層が0~100nmに抑えられたボルトを用い、雰囲気を酸素濃度が10ppm(体積基準)以下の不活性ガス雰囲気とする請求項6に記載のボルトの製造方法。
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