EP1956100B1 - Verfahren zum warmumformen eines stahlmaterials und dadurch erhaltenes stahlmaterial - Google Patents

Verfahren zum warmumformen eines stahlmaterials und dadurch erhaltenes stahlmaterial Download PDF

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EP1956100B1
EP1956100B1 EP06833094.3A EP06833094A EP1956100B1 EP 1956100 B1 EP1956100 B1 EP 1956100B1 EP 06833094 A EP06833094 A EP 06833094A EP 1956100 B1 EP1956100 B1 EP 1956100B1
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Prior art keywords
steel
less
working
steel material
warm
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French (fr)
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EP1956100A4 (de
EP1956100A1 (de
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Yuuji Kimura
Tadanobu Inoue
Kaneaki Tsuzaki
Kotobu Nagai
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National Institute for Materials Science
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National Institute for Materials Science
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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
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0231—Warm rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0421—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the working steps
    • C21D8/0431—Warm rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/06—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
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/10—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies
    • C21D8/105—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of tubular bodies of ferrous alloys
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/0075—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for rods of limited 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
    • C21D9/0093—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
    • 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
    • C22C38/02—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

Definitions

  • the present invention relates to a steel to be worked into various structures, components of cars, and the like for use. More particularly, it relates to a steel for warm working to be subjected to warm working, and a warm working method thereof, and a steel material and a steel component obtainable from the warm working method.
  • crystal grain refinement receives attention because of the following: it has both the effects of reduction of concentration of stress to the grain boundary and dilution at the grain boundary of impurity elements, and it can raise the brittleness fracture stress simultaneously with an increase in yield stress.
  • an attempt has been made to reduce the ferrite grain size as ultrafine as 1 ⁇ m or less in low carbon steel allowing for conservation of natural resources or recyclability, and thereby to achieve strengthening and the longer life of the steel.
  • Non-Patent Document 1-2 e.g., Non-Patent Document 1-2; Patent Document 1-2.
  • the crystal grains are required to be reduced in size to as ultrafine as 0.5 ⁇ m or less.
  • the reduction in grain size to as ultrafine as 0.5 ⁇ m or less is very difficult with a thermo-mechanical treatment of a steel intended for mass production.
  • ultrafine grains of 0.5 ⁇ m or less can be obtained with a super heavy deformation method such as MM (Non-Patent Document 3) or ARB (Non-Patent Document 4) of powder metallurgy.
  • MM Non-Patent Document 3
  • ARB Non-Patent Document 4
  • ultrafine grain steel shows almost no uniform elongation, and the elongation is mostly caused by ununiform deformation due to necking, resulting in a large reduction in ductility.
  • the same early plasticity instability as this is also observed in a pure iron wire dislocation strengthened by wire drawing (Non-Patent Document 5).
  • strengthening of a steel unfavorably largely reduces various characteristics such as ductility, toughness, delayed fracture resistance characteristics, fatigue characteristics, formability.
  • various characteristics such as ductility, toughness, delayed fracture resistance characteristics, fatigue characteristics, formability.
  • a very versatile low alloy martensitic steel is strengthened to 1.2 GPa or more, it is remarkably reduced in toughness, delayed fracture resistance characteristics, and the like. For this reason, the high strength steel is largely prevented from being put into practical use. Under such circumstances, there is a strong demand for simultaneously achieving higher strength and higher toughness, and the improvement of the delayed fracture resistance characteristics of a low alloy steel.
  • High temperature tempering is carried out at about 550°C or more, and A1 point or less. According to this, there are advantages as follows: (1) the internal stress introduced upon quenching can be largely reduced accompanied by recovery of dislocation; (2) the coherent precipitate (e.g., film-like cementite) reducing the fracture toughness can be made incoherent (be spheroidized), and other advantages. For this reason, for a steel for a mechanical structure particularly requiring toughness, tempering is generally carried out in the vicinity of 650°C. However, within such a temperature range, the dispersed second-phase particles also grow with ease during tempering, and hence the reduction in strength of the steel is inevitable.
  • the coherent precipitate e.g., film-like cementite
  • Patent Document 3 it is known (Patent Document 3) that refinement of austenite grains of several micrometers or less improves the toughness of the tempered martensitic steel (Non-Patent Document 13), and improves the delayed fracture characteristics (Patent Document 4).
  • Non-Patent Document 13 Non-Patent Document 13
  • Patent Document 4 improves the delayed fracture characteristics
  • dispersion of pinning grains effective for suppressing the growth of austenite reduces the austenitization temperature, austenitization for a rapid short time using high-frequency heating, and the like.
  • Ausforming is a treatment in which an austenitized steel is quenched to a metastable austenite range, and processed at the temperature, followed by quench hardening, thereby to cause martensite or bainite transformation, and then, tempering is carried out It has a feature of being capable of strengthening the steel without much impairing the toughness.
  • ausforming it is considered as follows. The effects such as (1) refinement of packets or blocks regarded as effective crystal grains, (2) succession of dislocation from worked austenite to martensite, and (3) pinning of dislocation by carbon atoms or carbides occur in an overlapping manner, thereby to strengthen the steel.
  • Non-Patent Documents 16 and 17 For enhancing the toughness of a steel, it is also effective to effect the formation of a fiber structure in the inside by cold or warm working. This has already been proposed for the steel subjected to an ausforming treatment (Non-Patent Documents 16 and 17), the heavy cold drawing high-strength low carbon wire rod (Patent Document 5), piano wire, pure iron wire (Non-Patent Document 5), and the like.
  • a manufacturing method of a formed product characterized by the following: a material having an ultrafine structure is warm worked or cold worked, and a steel material including drawn ferrite grains with a minor axis of 3 ⁇ m or less is used as a material; it is not subjected to a refining treatment, and only forming is carried out, and a refining treatment is not carried out (Patent Document 3).
  • Non-Patent Documents 11 and 12 wherein, for a steel having a duplex phase structure such as a tempered martensite structure, warm working is applied for the purpose of obtaining a worked structure before a quench hardening treatment for refining a reverse-transformed austenite.
  • the strength of the steel is achieved by a refining treatment following warm working. For this reason, an attempt has not been made to use the material in the as-warm worked state of the tempered martensite structure.
  • Non-Patent Document 18 For warm straightening working of a high carbon steel having a carbon content of 0.7 wt% or more, an over 1.8 GPa class wire rod can be obtained. However, the elongation of the wire rod is as low as around 6 % (Non-Patent Document 18).
  • Patent Document 7 describes toughening Mn-Mo or Mn-Cr-Mo steels by heating, tempering and quenching and warm working (stretching 24%)
  • prior-austenite grain refinement and ausforming are important toughness-enhancing technology of a steel, and studies and inventions thereon add up to massive amounts.
  • quench hardening and tempering are basic, so that strengthening receives restrictions by the problems of hardenability and quenching crack, and the problem of temper brittleness.
  • the amount of the dispersed second-phase particles of carbides or the like necessary for strengthening also increases, which makes softening by spheroidizing or the like difficult.
  • cracks occur in the inside of the material in the process of forming the material into a component by cold forging or the like, or other problems occur.
  • the studies and inventions regarding warm working up to this point mainly aim at forming into a member and manufacturing of the prior structure. For this reason, in most cases, a relatively soft matrix structure such as a ferrite or pearlite structure with a low deformation resistance, or a martensite structure subjected to tempering at high temperatures is used as a starting material. Thus, under such conditions as to result in reduction of the deformation resistance, warm working is carried out. Further, the fine duplex structure is not formed in consideration of the dispersion state of the dispersed second-phase particles and the thermal stability.
  • the present inventors conducted a close study in order to resolve the foregoing problems. As a result, they made the following inventions.
  • the softening resistance when the steel is heated namely, the thermal stability and the total amount of the matrix structure and the dispersed second-phase particles are controlled.
  • a particle dispersion type fibrous structure can be formed when the steel is subjected to warm working, and the Vickers hardness after warm working can be set at 3.7 ⁇ 10 2 or more.
  • the steel for warm working capable of being tremendously improved in toughness while keeping the tensile strength of 1.2 GPa or more at ordinary temperatures.
  • the structure of the steel for warm working as the prior working structure is transformed into an ultrafine duplex structure including dispersed second-phase particles such as carbide particles finely dispersed therein by using martensite transformation or bainite transformation.
  • martensite transformation or bainite transformation it becomes possible to effect the formation of a fiber structure even in the inside with efficiency when the steel is subjected to warm working.
  • the alloy composition excellent in cost efficiency and recyclability can achieve strengthening of the steel obtained when the steel is subjected to warm working.
  • the fourth invention it is possible to disperse dispersed second-phase particles which are finer and excellent in hydrogen trapping property. Further, it is possible to strengthen the steel material obtained when the steel is subjected to warm working, and to largely enhance the toughness in a low temperature range, and the delayed fracture resistance characteristics.
  • a fiber structure can be formed to obtain a high toughness.
  • the equipment warm working equipment which has been conventionally put into practical use can be utilized. Therefore, the invention has a very high practical utility.
  • a dense fiber structure with an average spacing of the minor axes of 1 ⁇ m or less is developed, and in accordance with the tenth invention, a dense fiber structure with an average spacing of the minor axes of 0.5 ⁇ m or less is developed.
  • steel materials which have been much more enhanced in strength, toughness, and workability than before warm working.
  • the average particle diameter of the major axes of the dispersed second-phase particles by controlling the average particle diameter of the major axes of the dispersed second-phase particles to 0.1 ⁇ m or less, it is possible to implement much more strengthening and enhancement of the toughness with dispersion of a small amount of the dispersed second-phase particles.
  • the steel material not only has a high toughness and tensile strength, but also has a secondary workability. For this reason, there are implemented a steel plate and a steel rod wire which have been tremendously enhanced in practical utility, usable for manufacturing various components and products.
  • a bolt excellent in impact strength and delayed fracture resistance in which a fiber structure is formed in the root of the thread of the screw part to which a stress particularly concentrates.
  • a high strength steel multiphased by fine dispersion of a small amount of dispersed second-phase particles is provided.
  • an ultrahigh strength steel which is hard to soften and is hard to form is applied with prescribed deformation in a temperature range in which the deformation resistance is reduced and no cracks occur in the material, to be formed into a prescribed shape (thin plate, thick plate, wire rod, or component).
  • a prescribed shape thin plate, thick plate, wire rod, or component.
  • conventional spheroidizing and quench hardening and tempering treatments after component forming are omitted.
  • an ultrafine duplex phase structure is developed into a fibrous form.
  • a high strength steel largely improved in the ductility, particularly the toughness, and the delayed fracture resistance characteristics in the relation of trade-off balance with high strength, and a member thereof.
  • a material satisfying given specific conditions can form a particle dispersion type fiber structure far more excellent in toughness and delayed fracture resistance characteristics even than conventional ausformed steels in the member. Namely, the pinning effect due to fine dispersion or precipitation of the second-phase particles is effectively used.
  • the material in a temperature range in which recovery of the dislocation introduced by deformation appropriately occurs, but primary recrystallization or remarkable grain growth does not occur, the material is deformed, and applied with prescribed strain, thereby to refine crystal grains.
  • a fiber structure having a further narrower crystal grain boundary spacing is developed. As a result, it is possible to suppress not only the occurrence of cracks but also the propagation of cracks, and thereby to largely enhance the fracture toughness.
  • Alloy elements high in carbide forming ability such as Mo, V, W, Ta, Ti, and Nb form nano-size alloy carbides such as Mo2C, V4C3, W2C, TaC, NbC, and TiC in a temperature range in the vicinity of 500°C to 600°C independently from already existing cementite. For this reason, addition of these alloy elements is effective for strengthening of the steel.
  • the maximum value of precipitation strengthening due to these nano-size alloy carbides is obtained in the transition range of the strengthening mechanism of from Cutting to Orowan mechanism. However, at such an aging stage, much coherent strain occurs around the precipitates, so that the toughness of the steel is reduced.
  • the steel is tempered to the sufficiently overaged state of these carbides even somewhat sacrificing the strength of the steel.
  • the dynamic precipitation of these alloy carbides due to warm working it is also possible to effect incoherent precipitation of the carbides without causing much growth of the carbides even in the precipitation transition temperature range. Namely, it is also possible to make the maximum use of precipitation strengthening of the alloy carbides due to the Orowan mechanism. Further, the same effects can also be expected for precipitation of intermetallic compounds including the alloy elements, and Ni, Al, or the like, nitrides, oxides, Cu particles, and the like.
  • the steel for warm working of the invention changes in the dispersion state of the dispersed second-phase particles and the matrix structure during warm working to be performed thereon. Therefore, it is configured such that the lower limit of the equation (2) is set with respect to the hardness (structure) of the non-worked material obtained from the heat treatment simulating the heat history of warm working. Namely, as described below, the structure state is expressed by the hardness.
  • dispersion (precipitation) strengthening by the dispersed second-phase particles depends upon the dispersion conditions such as the volume fraction of the secondary phase duplex particles, and the size, hardness, and shape of the particles.
  • dispersion strengthening is caused by the Orowan mechanism, from the following equation (A) ( TEKKOU NO SEKISYUTSUSEIGYO METARAJII SAIZENNSENN, (the Iron and Steel Institute of Japan) (2001) P. 69 )
  • the amount of dispersion strengthening increases with a decrease in particle diameter (d) and with an increase in volume fraction (f).
  • the dispersion conditions (and the dispersive power) of the dispersed second-phase particles have close relation with the hardness.
  • the dislocation ceases to be pinned by the particles.
  • the particles come to be sheared by dislocation, so that the Orowan mechanism ceases to hold.
  • the so-called Cutting mechanism in which particles are sheared by dislocation, the amount of dispersion strengthening increases with an increase in particle diameter. Namely, the minimum particle diameter with which the Orowan mechanism holds can provide the maximum amount of dispersion strengthening.
  • the minimum particle diameter capable of achieving the maximum dispersion strengthening depends upon the hardness of the particles, and decreases in inverse proportion to the hardness of the particles ( TEKKOU NO SEKISYUTSUSEIGYO METARAJII SAIZENNSENN, (the Iron and Steel Institute of Japan) (2001) P. 69 ). Therefore, when comparison is made in terms of the same volume fraction, the minimum particle diameter with which the Orowan mechanism holds decreases with an increase in hardness of the particles. Accordingly, the maximum amount of particle dispersion strengthening also increases.
  • TiC is capable of carrying out effective dispersion particle strengthening because of its higher hardness and smaller density among alloy carbides.
  • a particle dispersion strengthening amount of about 0.9 GPa (TS (GPa) is nearly equal to 0.0032 HV, HV 2.8 ⁇ 10 2 ) is expectable in dispersion at a volume fraction of 7 ⁇ 10 -3 .
  • Ti and C necessary for precipitating TiC at a volume fraction of 7 ⁇ 10 -3 are in an amount of 0.35 wt% and 0.087 wt%, respectively.
  • the strength of the matrix of the practical ferrite steel is about 0.3 GPa (about 0.9 ⁇ 10 2 in HV). Therefore, the room temperature strength of the steel including the TiC dispersed in the ferrite matrix is expected to be 1.2 GPa or more (HV 3.7 ⁇ 10 2 or more).
  • a size of 7 nm or more can sufficiently satisfy an HV of 3.7 ⁇ 10 2 only by dispersion strengthening at a volume fraction as small as 7 ⁇ 10 -3 .
  • the same effects as with this are also expectable for the dispersed second-phase particles including a carbonitride, an intermetallic compound, an oxide, Cu particles, and the like.
  • dispersed second-phase particles for example, specifically, there can be considered carbides such as Mo2C, V4C3, W2C, TaC, NbC, and TiC, oxides such as Fe3O4, Fe2O3, Al2O3, Cr2O3, SiO2, and Ti2O3, nitrides such as AlN, CrN, and TiN, intermetallic compounds such as Ni3Ti, NiAl, TiB, Fe2Mo, Ni3Nb, and Ni3Mo, metal particles such as Cu particles, and the like.
  • carbides such as Mo2C, V4C3, W2C, TaC, NbC, and TiC
  • oxides such as Fe3O4, Fe2O3, Al2O3, Cr2O3, SiO2, and Ti2O3, nitrides such as AlN, CrN, and TiN
  • intermetallic compounds such as Ni3Ti, NiAl, TiB, Fe2Mo, Ni3Nb, and Ni3Mo
  • metal particles such as Cu particles, and the like
  • metal carbide particles of Mo, Ti, or the like generally have a size of around 10 nm, and can effectively effect strengthening even by dispersion in an amount as small as less than 10 ⁇ 10 -3 in volume fraction.
  • the size of the dispersed second-phase particles and the distribution in the matrix structure also varies according to segregation of alloy elements or the like, or other factors. Therefore, in the invention, considering such that a fine crystal structure can be obtained with stability by warm working even when there are variations in distribution of the dispersed second-phase particles, the volume fraction at room temperature of the dispersed second-phase particles is specified at 7 ⁇ 10 -3 or more.
  • the volume fraction of the dispersed second-phase particles is preferably set at 20 ⁇ 10 -3 or more.
  • the upper limit of the volume fraction of the dispersed second-phase particles has no particular restriction in effecting strengthening.
  • it is preferably set at 12 ⁇ 10 -2 or less.
  • particle dispersion strengthening by the Orowan mechanism is expected to become remarkable in the region of several tens nanometers or less from the equation (A).
  • the average particle diameter of the dispersed second-phase particles is desirably 0.5 ⁇ m or less, and more preferably 0.1 ⁇ m or less as the steel for warm working.
  • the foregoing conditions are predicated upon the fact that the dispersed second-phase particles do not grow even in the temperature range equal to or higher than that of the third or more stage of tempering of 350°C.
  • the steel in order for the steel to have a strength of 1.2 GPa or more even after warm working, it becomes a necessary condition that, during heating and working, and after working, in addition to the matrix structure, particularly, the dispersed second-phase particles do not undergo Ostwald ripening, resulting in a reduction of the strength.
  • the wording "in a prescribed temperature range” denotes that the foregoing conditions may be satisfied at any temperature of from 350°C to the Ac1 point, and means that the foregoing conditions are not required to be satisfied over the whole temperature range.
  • the material undergoes remarkable aging hardening or secondary hardening to have a hardness of H or higher only in a given temperature range within the foregoing range, it can serve as the steel for warm working of the invention.
  • a fibrous structure with an average grain size of 3 ⁇ m or less is obtained by the two effects of refinement of the matrix structure by impartment of a predetermined strain and grain boundary pinning by TiC.
  • the necessary and sufficient condition of the prior working structure is as follows: the lower limit value of the volume fraction of the dispersed second-phase particles is set at 7 ⁇ 10 -3 , and the steel after any heat treatment of annealing, tempering, and aging under the condition T(logt + 20) ⁇ 1.4 ⁇ 10 4 has a hardness of HV ⁇ (5.2 - 1.2 ⁇ 10 -4 ⁇ ) ⁇ 10 2 .
  • the invention has the following features as a steel for warm working: fine dispersion or precipitation of the dispersed second-phase particles in the matrix structure as particle dispersion strengthening particles, and the structure control to enhance the thermal stability of the dispersed second-phase particles.
  • the dispersion conditions of the dispersed second-phase particles and the matrix structure variously change. Therefore, although not limited by the room-temperature structure form, all the steels with a strength of 1.2 GPa or more except for the steels having a pearlite structure as the main structure can be considered as the steel for warm working.
  • martensitic steels tempered martensite structures
  • the second steel for warm working of the invention is configured such that 80 percent by volume or more of the matrix structure is any single structure of martensite and bainite or a mixed structure thereof.
  • the width of the block regarded as the effective crystal grain of martensite is 1 ⁇ m or less in a medium carbon low alloy steel ( Scripta Mater., 49 (2003), P. 1157 ).
  • the bainite structure also has a needle-like or plate-like structure form including a carbide finely dispersed therein.
  • a martensite or bainite structure having a temper softening resistance equal to or higher than that of a tempered martensitic steel of JIS-SCM430 steel is included in a volume of 80 % or more.
  • 20 % by volume or less structure other than martensite or bainite and a mixed structure thereof may be accounted for by any structure such as a ferrite, pearlite, or austenite structure. This is for the following reason.
  • Such a ferrite, pearlite, or austenite structure, or the like decomposes or disappears, or changes into a microstructure during a warm thermo-mechanical treatment. Therefore, when it is present in an amount of 20 % by volume or less, it is judged as no problem.
  • the third to fifth steels for warm working of the invention are alloy designed based on the foregoing findings.
  • the gist resides in a steel for warm working characterized by containing, as the chemical composition, C: 0.70 wt% or less, Si: 0.05 wt% or more, Mn: 0.05 wt% or more, Cr: 0.01 wt% or more, Al:0.5 wt% or less, O: 0.3 wt or less, and N: 0.3 wt% or less, and the balance substantially being Fe and inevitable impurities: Whereas, the following and the like can be considered.
  • the steel for warm working further contains one or two or more selected from a group consisting of Mo: 5.0 wt% or less, W: 5.0 wt% or less, V: 5.0 wt% or less, Ti: 3.0 wt% or less, Nb: 1.0 wt% or less, and Ta: 1.0 wt% or less, or contains one or two or more of Ni: 0.05 wt% or more and Cu: 2.0 wt% or less.
  • C forms a carbide particle, and is the most effective component for strengthening.
  • the content exceeds 0.70 wt%, the toughness degradation is caused. For this reason, the content is set at 0.70 wt% or less.
  • the content is preferably 0.08 wt% or more, and more preferably 0.15 wt% or more.
  • Si is an effective element for enhancing the strength of the steel by deoxidation and solid solution in ferrite, and finely dispersing cementite. Therefore, the content is set at 0.05 wt% or more inclusive of the one added as a deoxidizer, and to remain in the steel. The upper limit is not particularly restricted for strengthening. However, in view of the workability of the steel material, the content is preferably set at 2.5 wt% or less.
  • Mn is an effective element for reducing the austenitization temperature, and refining austenite. In addition, it is an effective element for the hardenability, and being dissolved in a solid solution form in cementite and suppressing coarsening of cementite.
  • the content is set at 0.05 wt% or more. More preferably, the content is 0.2 wt% or more.
  • the upper limit is not particularly restricted for strengthening. However, in view of the toughness of the resulting steel material, the content is preferably set at 3.0 wt% or less.
  • Cr is an effective element for improving the hardenability, and it is an element having a strong action of being dissolved in a solid solution form in cementite and delaying the growth of cementite. Further, it is also one of the important elements in the present invention for forming a high Cr carbide which is thermally more stable than cementite, or improving the corrosion resistance by being added in a relatively larger amount. Therefore, Cr is required to be contained at least in an amount of 0.01 wt% or more. It is contained in an amount of preferably 0.1 wt% or more, and more preferably 0.8 wt% or more.
  • Al is an effective element for deoxidization and forming an intermetallic compound with an element such as Ni and enhancing the strength of the steel. However, excessive addition reduces the toughness. Therefore, the content is set at 0.5 wt% or less.
  • the content is preferably set at 0.02 wt% or less, and further restrictively 0.01 wt% or less.
  • O oxygen
  • oxygen effectively acts not as an inclusion but as grain growth preventing or dispersion strengthening particles when it can be finely and uniformly dispersed as an oxide.
  • the content is set at 0.3 wt% or less.
  • the oxide is not used as the dispersion strengthening particles, the content is preferably set at 0.01 wt% or less.
  • N nitrogen
  • nitrogen effectively acts as grain growth preventing or dispersion strengthening particles when it can be finely and uniformly dispersed as a nitride.
  • the content is set at 0.3 wt% or less.
  • the content is preferably set at 0.01 wt% or less.
  • Mo is an effective element for strengthening the steel in the invention. It not only improves the hardenability of the steel, but also is dissolved in a small amount in a solid solution form also in cementite to make cementite thermally stable. Particularly, completely separately from cementite, it newly causes separate nucleation of an alloy carbide on dislocation in the matrix phase, thereby to effect secondary hardening, resulting in strengthening of the steel. Further, the formed alloy carbide is effective for grain refinement, and is also effective for replacement of hydrogen. Therefore, Mo is contained in an amount of preferably 0.1 wt% or more, and more preferably 0.5 wt% or more.
  • the upper limit of the content is set at 5 wt%. From the viewpoint of cost efficiency, the content is preferably set at 2 wt% or less.
  • Ni is an element effective for improving the hardenability, and effective for reducing the austenitization temperature, and refining austenite, improving the toughness, and improving the corrosion resistance. Further, it is also an effective element for forming an intermetallic compound with Ti or Al, and precipitation strengthening the steel when it is contained in a proper amount. When the content is less than 0.01 wt%, desired effects cannot be obtained. In the present invention, the content is set at 0.05 wt% or more. Ni is more preferably contained in an amount of 0.2 wt% or more. The upper limit is not particularly restricted. However, Ni is an expensive element, and hence it is preferably contained in an amount of 9 wt% or less.
  • Cu is a detrimental element causing hot brittleness. But on the other hand, when it is added in a proper amount, it causes precipitation of fine Cu particles at 500°C to 600°C, thereby to strengthen the steel. When it is added in a large amount, it causes hot brittleness. Therefore, the content is set at 2 wt% or less which is roughly the maximum amount of solid solution into ferrite.
  • P phosphorus
  • S sulfur
  • P or S reduces the grain boundary strength, and hence it is an element which is desired to be removed as much as possible.
  • Each content is preferably set at 0.03 wt% or less.
  • various ones can be considered according to the methods for manufacturing martensite structures or bainite structures of JIS standard, and the like. These are not limited to dissolution and forging methods.
  • other manufacturing methods such as powder metallurgy can also be used.
  • the following procedure or the like is also possible.
  • a technique such as a ball milling method, most undissolved compounds such as oxides in the steel are steel powder dispersed with a size of nanometer size, and then, ( ISIJ International, 39 (1999), p176 ), such a mechanical milled powder is consolidated and formed in a proper temperature range to obtain an objective bulk body.
  • the warm working method of the invention is characterized by subjecting any steel for warm working described above to warm working for applying 0.7 or more strain in the temperature range of 350°C or more and Ac1 point -20°C or less. It can also be considered that after performing warm working, an aging treatment is performed in the temperature range of 350°C or more and Ac1 point or less. According to such warm working, the following advantages can be obtained:
  • a working temperature more specifically, for example, in the case of a medium carbon low alloy steel for use as a steel for general mechanical structures, including a martensite structure as the matrix, it can be set at 350°C or more roughly corresponding to the third stage of tempering in which cementite precipitates.
  • a working temperature in order to effectively use an alloy carbide, an intermetallic compound, Cu, or the like as the dispersed second-phase particles, it is desirable that working is carried out in the temperature range of 500°C to 650°C which is the precipitation temperature of the second-phase particles.
  • phase transformation such as pearlite transformation or martensite transformation is effected during the cooling process.
  • phase transformation such as pearlite transformation or martensite transformation is effected during the cooling process.
  • the upper limit temperature of working is set at Ac1 point -20°C.
  • the combination of the working temperature and time of the material when the hardness is arranged by a tempering parameter ⁇ , the combination such that the Vickers hardness at room temperature when the material is subjected to any of annealing, tempering, and aging treatments in the as-unworked state does not become 3.7 ⁇ 10 2 or less in HV is preferable in order to obtain a strength of 1.2 GPa or more after warm working.
  • the time required for working is required to be shorten in view of the softening resistance and the heating time of the material.
  • the degree to which the structure is developed depends upon the prior working structure, the working temperature, and the strain amount In other words, the necessary strain amount varies according to the prior working structure or the working temperature. Therefore, although the strain amount cannot be strictly specified, it is preferable that a strain of 0.7 or more, and more preferably 1 or more is imposed when a fibrous structure is desired to be formed in the inside of the material.
  • a strain amount is preferably 1 or more, and more preferably 1.5 or more.
  • the strain to be imparted may be introduced not only in a single working pass, but in a plurality of divided passes. Further, the direction of working is not constantly limited to the same direction. Still further, the inter-pass time is also not particularly restricted. Furthermore, the process also includes imposing a prescribed strain not over the entire region of the material to be worked, but on a specific region (e.g., the surface layer requiring strengthening, or R part of a component).
  • the actual strain amount can be understood only after considering the material characteristics of the material to be worked, the friction conditions (e.g., the type or the presence or absence of a lubricant) of the roll (mold for forging) and the material to be worked, the deformation of the roll (mold for forging), rolling (forging) rate, the rolling (forging) temperature, and the like.
  • the friction conditions e.g., the type or the presence or absence of a lubricant
  • the cumulative rolling reduction when the cumulative rolling reduction is 45 % or more for plate rolling intended for the plane strain state, or when the cumulative reduction of area is 45 % or more for wire rod rolling, it can be considered that a strain of 0.7 or more has been introduced into the entire region of the material to be worked.
  • the cumulative rolling reduction or the cumulative reduction of area when the cumulative rolling reduction or the cumulative reduction of area is 58 % or more, it can be considered that a strain of 1 or more has been introduced into the entire region of the material to be worked.
  • a strain of 0.7 or more may be introduced into the entire region or into a specific region of the material to be worked under the influence of friction or the like. Therefore, in that case, it is necessary to quantitatively study the amount of introduced strain by numerical analysis.
  • the steel material of the invention is a steel obtained by warm working a steel for warm working as described above. It is characterized in the following respects: it has a matrix structure including a fibrous crystal having an average grains size of the minor axis of 3 ⁇ m or less; the second-phase particles are finely dispersed in the matrix structure at a volume fraction of 7 ⁇ 10 -3 or more at room temperature; and the Vickers hardness at room temperature is 3.7 ⁇ 10 2 or more in HV.
  • the matrix structure in the steel material of the invention includes a fibrous ferrite crystal with an expansion degree (aspect ratio) of more than 2, and typically with an aspect ratio of 5 or more, in which the second-phase particles are finely dispersed.
  • the upper limit of the average spacing (i.e., minor axis average grain diameter) of the matrix structure including the fibrous crystal is set at 3 ⁇ m.
  • crystal grains denotes the crystal grains surrounded by the grain boundaries with a crystal orientation difference of 15° or more.
  • the average particle diameter of the major axis of the dispersed second-phase particles is larger than 0.3 ⁇ m, particle dispersion strengthening can be hardly expected.
  • the average particle diameter of the major axes is desirably 0.3 ⁇ m or less.
  • the effects of crystal grain refinement becomes especially remarkable in the region in which the average crystal grain diameter is 1 ⁇ m or less; and particle dispersion strengthening due to the Orowan mechanism, in the region in which the average particles diameter is 0.1 ⁇ m or less.
  • the minor axis average grain diameter of the fibrous crystal is set at 1 ⁇ m or less, and further 0.5 ⁇ m or less.
  • the average particle diameter of the major axes of the dispersed second-phase particles is also set at 0.1 ⁇ m or less, and further 0.05 ⁇ m or less according to the refinement of the matrix structure.
  • strengthening mechanisms such as solid solution strengthening and dislocation strengthening can also be applied.
  • the effects of superposition of these strengthening mechanisms lead to provide such a high performance material as unpredictable with simple addition of the strengthening mechanisms.
  • Such fine fiber structures can be formed by warm forming of rod wire materials, screw members of bolts, and the like, including plate materials. Particularly even when the cumulative strain amount is small, a fiber structure can be formed in the surface layer part which has locally undergone intense deformation or the like. This can largely improve the characteristics of various components and the desirable parts.
  • Table 1 shows the steel components (A to K, M, N, and O) within the scope of the invention and the steel component (L) outside the scope.
  • carbides are used as the dispersed second-phase particles.
  • Table 2 shows the volume fractions of alloy carbides dispersible as the dispersed second-phase particles and cementites for the steels of the compositions of Table 1.
  • the steels of the examples cover martensitic steels of from SCM435 to 2 GPa class secondary hardened steels except for Co-added maraging steels.
  • Carbides of various stoichiometric compositions are present in actual steels according to the components of the steel and the heat treatment conditions. For this reason, it is difficult and not practical to strictly measure the volume fraction of the dispersed second-phase particles by chemical analysis or structure observation. Under such circumstances, the present inventors determined the volume fraction of each carbide from the known theoretical density determined by the structural analysis or the like of the carbide ( KAGAKU DAIJITENN, TOKYO KAGAKU DOJIN Co. Ltd., (1989), P. 1361 to 1363 ). The approximate expressions and the like are as shown in Table 3.
  • the alloy elements respectively combine with carbons in order of decreasing carbide forming ability (Nb > Mo > Cr > Fe, and the like) to form carbides.
  • Nb or Mo it is well known that it is an element which tends to form its specific carbide in the steel and is less likely to be dissolved in cementite.
  • precipitation of NbC or Mo2C is assumed.
  • Mo in an amount of 0.002 wt% can be sufficiently dissolved in a solid solution form in cementite, and hence it is excluded from the estimation of the volume fraction of the Mo carbide.
  • the amount of carbide serving as the dispersion strengthening particles to be dispersed depends upon the carbon content in a medium carbon low alloy steel. Particularly, when there is no possibility that a alloy carbide with a sufficiently large density with respect to cementite is formed, or when the amount of elements for forming alloy carbides to be added is small, the amount of the second-phase particles to be dispersed is roughly determined by the amount of cementite. Namely, as shown in Table 2, in the steels having a C content of 0.2 wt% or more used in examples of Table 1, the total amount of the volume fractions of the second phase sufficiently exceeds 7 ⁇ 10 -3 .
  • thermo-mechanical treatment applied in examples basically includes (1) a solution treatment and working for reducing coarse undissolved carbides; (2) a quench hardening treatment and tempering for obtaining a tempered martensite or bainite structure as a structure of the steel for warm working of the invention; and (3) warm working also serving as shape forming into a component.
  • a solution treatment and working for reducing coarse undissolved carbides (2) a quench hardening treatment and tempering for obtaining a tempered martensite or bainite structure as a structure of the steel for warm working of the invention; and (3) warm working also serving as shape forming into a component.
  • FIG. 3 shows a quench hardening process from the worked austenite (elongated austenite) structure by an ausforming treatment in a metastable austenite range.
  • a square bar of about 40 mm square ⁇ 120 mm in length cut from a hot rolled steel sheet or a forged material was subjected to the steps up to the quench hardening treatment in the thermo-mechanical treatment patterns 1, 2, and 3, thereby to obtain a martensite single structure close to nearly 100 % by volume.
  • the square bar was heated to a prescribed temperature for 0.5 hour, and tempered. Then, it was subjected to warm rolling working to a prescribed reduction of area by the use of a groove roll, and applied with a strain, and air cooled.
  • the cross section in parallel with the rolling direction (RD) was polished and observed by the use of an optical microscope, a transmission electron microscope (TEM), and a FE-SEM with EBSP analyzer.
  • the polished surface was corroded with picric acid alcohol, and the prior-austenite grain boundary was revealed.
  • the prior-austenite grain diameter was determined according to the comparison method or the cutting method specified in JIS G 0552.
  • the average particle diameter of the dispersed second-phase particles was determined in the following manner. By the use of TEM or SEM, 3 or more visual fields were observed at a magnification of 10000 times to 100000 times to measure the length of each major axis of a total of 250 or more particles.
  • the maximum particle diameter is allowed to correspond to the length of the major axis of the largest carbide of the carbides measured.
  • the average grain diameters of the minor axes and the major axes of the elongated grains in the fiber structure according to EBSP analysis, the average section lengths of the minor axes and the major axes of the elongated crystal grains having a crystal orientation difference of 15° or more were measured with a cutting method (see FIG. 5 ).
  • the hardness of the resulting steel material was measured under a load of 20 kg and for a holding time of 15 s by means of a Vickers hardness tester according to the testing method specified in JIS Z 2244.
  • the tensile test was performed at ordinary temperatures by means of an Instron type tensile test machine according to the testing method specified in JIS Z 2241, for 1) a JIS No. 14 proportional test piece with a parallel part diameter of 3.5 mm, a length of 24.5 mm, and a mark-to-mark distance of 17.5 mm, or with a size of 6 mm, a length of 42 mm, and a mark-to-mark distance of 30 mm, or for 2) a JIS No. 4 sub-size test piece with a parallel part diameter of 10 mm, a length of 45 mm, and a mark-to-mark distance of 35 mm.
  • the cross head speeds were 0.5 mm/min and 10 mm/min for 1) JIS No. 14A and 2) JIS No. 4, respectively.
  • the elongation was measured until rupture by mounting an extensometer on each test piece.
  • the impact test was performed according to the testing method specified in JIS Z 2242 for a U notch or V notch test piece of 55 mm in length and 10 mm in height and width manufactured by cutting machining from a steel material of 1.8 cm 2 or more in cross section.
  • the hydrogen embrittlement characteristics were evaluated at room temperature at a cross head speed of 0.005 mm/min using a slow strain rate tensile test machine for each notch test piece of 10 mm in diameter, 6 mm in notch bottom diameter, and 4.9 in stress concentration coefficient.
  • the test was carried out after setting the following conditions.
  • the average hydrogen amount in the test piece was changed by 72-hour cathode charge with varying charged solution and current density, to perform Cd plating. This prevents hydrogen in the test piece from dissipating.
  • the analysis of hydrogen was carried out by a thermal desorption analysis using a quadrupole mass spectrometer for a sample from which Cd-plating had been removed.
  • the hydrogen to be released up to 300°C was defined as diffusive hydrogen, and determined.
  • Table 4 summarizes the manufacturing conditions and the structure form of each steel for warm working, and the quench hardening and tempering conditions of the unworked material and the hardness thereof, and the results of evaluation of the suitability as the steel for warm working of the invention.
  • Example 1 M 1 12.0 0 12.0 - - 0.5 8.5 1.55 5.4 For working 13 1.75 4.4 O 2 12.0 40 9.0 - - 0.06 4.8 1.55 4.2 For working 14 3 12.0 40 10.2 0 5.8 0.06 4.6 1.55 4.2 For working 15 33 5.8 . 0.05* 5.5 1.55 4.7 For working 16 55 5.7 0.04* 5.6 1.55 4.9 For working 17 70 5.8 0.03* 5.7 1.55 5.0 For working 18 1) Thermo-mechanical treatment pattern * For ausformed (elongated grain) structure, the grain diameter of the minor axis is measured.
  • the volume fraction of cementite is 33 ⁇ 10 -3 .
  • FIG. 5 shows an example of analysis of the structure of the material obtained by subjecting an I steel to ⁇ transformation at 11.0 ⁇ 10 2 °C, followed by water cooling, and subjecting it to a tempering treatment at 5.0 ⁇ 10 2 °C for 1.5 hours, followed by warm groove rolling.
  • the cumulative strain amount imparted at this step is 2.4
  • the hardness is HV 3.7 ⁇ 10 2 .
  • the average grain diameter of the minor axes of crystal grains having a crystal orientation difference of 15° or more was measured with a cutting method.
  • the average grain diameter of the minor axes of the elongated crystal grains was found to be 0.3 ⁇ m.
  • the fiber structure has been developed in a complicated form, so that it was not possible to measure the average grain diameter of the major axes.
  • the 287 carbide particle diameters (major axis lengths) were measured by TEM. As a result, the average particle diameter of the carbide was found to be 0.06 ⁇ m, and the maximum diameter thereof was found to be 0.2 ⁇ m (d).
  • the inverse pole figure with respect to the rolling direction (RD) indicates that there is formed a fiber structure in which ⁇ 011>//RD texture has developed. Incidentally, also for other developed steels, the same textures were formed.
  • the cleavage plane of the Bcc iron is ⁇ 100 ⁇ . Therefore, the formation of such a ⁇ 011> fiber structure is considered to be very effective for fracture due to tensile deformation in the direction of fiber axis, flexural deformation receiving flexural moment along the direction of fiber, or the like.
  • Table 5 shows the relationship between the warm working conditions and the structure and hardness of the resulting warm worked material.
  • T and t in the table denote the working temperature and the working treatment time shown in FIGS. 1 to 3 , respectively.
  • Warm working conditions and structure form of worked material Steel type Warm working conditions Structure form of warm worked material 1)
  • Cumulative equivalent strain amount ⁇ *1 Average carbide particle diameter ( ⁇ m) *2 Maximum carbide particle diameter ( ⁇ m) *2 Minor axis average grain diameter of elongated grains
  • L1 ( ⁇ m) Major axis average grain diameter of elongated grains
  • L2 ( ⁇ m) Elongation rate (L2/L1)
  • Worked material hardness (HV) ⁇ 10 -2 For working 1 A 2 5.0 1.55 76 1.7 ⁇ 0.1 0.07 0.4 Immeasurable For working 2 B 1 5.0 1.55 76 1.7
  • Example 2 5.0 1.55 27 0.4 ⁇ 0.1 - - Immeasurable - 5.4
  • Example 4# 5.0 1.55 51 0.8 ⁇ 0.1 - - Immeasurable - 5.2
  • Example 5 6.0 1.75 76 1.7 ⁇ 0.1 0.3 0.3 Immeasurable - 4.6
  • Example 6 7.0 1.96 76 1.7 0.2 0.85 0.5 - - - 3.3 Comp.
  • Example 3 For working 4 C 2 5.0 1.55 76 1.7 ⁇ 0.1 0.15 0.3 Immeasurable - - 3.3 6.0
  • Example 7 6.0 1.75 76 1.7 ⁇ 0.1 - 0.4 Immeasurable - 4.8
  • Example 8 7.0 1.96 76 1.7 0.2 1.2 0.9 5.4 6 3.6 Comp.
  • Example 4 For working 5 D 1 5.0 1.55 76 1.7 ⁇ 0.1 0.4 0.3 Immeasurable - 4.6
  • Example 9 For working 6 E 1 5.0 1.55 76 1.7 ⁇ 0.1 0.3 0.3 Immeasurable - 4.6
  • Example 10 For working 7 F 1 5.0 1.55 76 1.7 ⁇ 0.1 0.3 0.3 Immeasurable - 4.9
  • Example 11 For working 8 G 1 5.0 1.55 76 1.7 ⁇ 0.1 0.4 0.4 Immeasurable - - 4.4
  • Example 12 6.0 1.75 76 1.7 0.2 0.6 0.5 Immeasurable - 3.8
  • Example 13 7.0 1.95 76 1.7 0.3 0.8 0.7 1.6 2
  • H For working 9 H 1 5.0 1.55 76 1.7 ⁇ 0.1 0.25 0.7 0.5 Immeasurable - 3.2 3.8 Comp.
  • Example 5 Example 14 5.0 1.55 57 1.0 ⁇ 0.1 0.3 1.4 Immeasurable - 3.8
  • Example 15 For working g 10 I 1 5.0 1.56 87 2.4 0.06 0.2 0.3 Immeasurable - 3.8 3.7
  • Example 16 6.0 1.76 87 2.4 0.08 0.25 0.4 Immeasurable - 3.1 Comp.
  • Example 6 7.0 1.96 87 2.4 0.11 0.48 0.7 2.8 4 2.4 Comp.
  • Example 7 For working 11 J 2 5.0 1.55 76 1.7 ⁇ 0.1 0.2 0.3 Immeasurable - 5.3
  • Example 17 For working 12 K 2 5.0 1.55 76 1.7 ⁇ 0.1 0.25 0.3 Immeasurable - 4.9
  • Example 18 For working 13 M 1 5.0 1.56 76 1.7 ⁇ 0.1 - 0.3 Immeasurable - 5.7
  • Example 19 6.0 1.76 76 1.7 ⁇ 0.1 - 0.7 Immeasurable - 4.3
  • Example 20 For working 14 O 2 5.0 1.55 76 1.7 ⁇ 0.1 ⁇ 0.1 0.9 Immeasurable - 5.6
  • Example 21 For working 15 3 5.0 1.55 79 1.8 ⁇ 0.1 ⁇ 0.1 0.8 Immeasurable - 5.7
  • Example 22 For working 16 5.0 1.55 68 1.3 ⁇ 0.1 ⁇ 0.1 1.0 Immeasurable - 5.9
  • the hardness of the worked material largely depends on the temper softening resistance.
  • a steel having a larger temper softening resistance provides a worked material with a higher hardness.
  • the matrix structure has been refined to as ultrafine as 0.5 ⁇ m or less in average width.
  • very fine particles are densely dispersed. For this reason, it was not possible to strictly determine the average carbide particle diameter.
  • comparison was made with the one including relatively large particles such as the I steel of FIGS. 5 it was possible to judge the diameter as less than 0.1 ⁇ m.
  • Tables 6 and 7 summarize Examples and Comparative Examples for the mechanical properties.
  • UE and VE denote the absorption energies of the U notch and V notch test pieces, respectively.
  • Example 4 in Table 6 is comparative.
  • [Table 6] Tensile deformation characteristics at room temperature of warm worked steel and impact absorption energy (J) at each test temperature Steel type Proof stress (GPa) Tensile strength TS (GPa) Uniform elongation (%) Total elongation (%) Reduction of area (%) TS ⁇ total elongation balance Room temperature toughness UE20 (J) Low temperature toughness UE-20 (J) Low temperature toughness UE-60 (J) Room temperature toughness VE20 (J) Low temperatur e toughness VE-20 (J) Low temperature toughness VE-60 (J) Ex.
  • the steels of the compositions herein shown have been subjected to proper alloy design and heat treatments so that the second-phase particles are finely dispersed as the steels for warm working.
  • Even unworked steels of Comparative Examples show a tensile strength ⁇ total elongation balance of 16 or more.
  • the developed steel subjected to warm working provides a larger tensile strength ⁇ total elongation balance than Comparative Examples.
  • even addition of carbon in an amount of about 0.2 wt% provides an over 1.5 GPa class almost the same as the quench hardening hardness when alloy elements such as Mo are properly added. Further, the ductility is excellent. These receive attention (A, D, and E steels). Further, O steel provides a 2 GPa class ultrahigh strength steel.
  • the impact absorption energy indicates that each developed steel has a far more excellent toughness than conventional high strength steels to a low temperature range.
  • FIG. 6 summarizes the relationship between the tensile strength and the impact value (U notch test piece) at room temperature.
  • the diagram there is also shown data of the steel for mechanical structure specified in JIS ( SHINNNIHON CYUUTANZOU KYOUKAI: GENNBAYOU KIKAIKOUZOUYOU HAGANEZAIRYOU DATA SHEETS (1995 )).
  • the impact value is largely reduced in the strength range of 1.2 GPa or more, and 70 J/cm 2 or less at a strength of 1.5 GPa or more.
  • the steels of the invention show a very high impact value of 150 J/cm 2 or more particularly even at a strength of 1.5 GPa or more.
  • FIG. 7 summarizes the relationship between the tensile strength and the absorption energy at room temperature (V notch test piece).
  • V notch test piece the absorption energy at room temperature
  • FIG. 7 summarizes the relationship between the tensile strength and the absorption energy at room temperature (V notch test piece).
  • JIS Institute of Metal Material Technology Fatigue Data sheet Reference 5
  • the invention is also superior in toughness in a high strength region to conventional ausformed steels, fine grain steels, maraging steels, and the like.
  • Fig. 8 shows the relationship between the test temperature and the absorption energy.
  • the absorption energy in the vicinity of room temperature is not only higher than that of the comparative steels, but also shows such a specific temperature dependency as to show the maximum value in a low temperature range and decrease.
  • the absorption energy in the vicinity of room temperature is not only higher than that of the comparative steels, but also shows such a specific temperature dependency as to show the maximum value in a low temperature range and decrease.
  • the absorption energy in the vicinity of room temperature is not only higher than that of the comparative steels, but also shows such a specific temperature dependency as to show the maximum value in a low temperature range and decrease.
  • the A steel of Example 1, or the B steel of Example 3 a peak is observed in the vicinity of -40°C; for the F steel of Example 11, in the vicinity of -100°C.
  • the maximum absorption energy at ordinary temperatures is about 90 J at a strength level of 1.5 GPa (Non-patent Document 17). Therefore, the following are the findings unprecedented and worthy of special mention.
  • the absorption energy in the vicinity of normal temperatures is far higher than that of existing ausformed steels, and in addition, the absorption energy shows the maximum value in a low temperature range of -40°C or less.
  • the excellent mechanical characteristics, particularly the high impact characteristics of the developed steels as described up to this point are largely due to the ultrafinefiber structure with ⁇ 011>/ /RD texture which has been densely developed by warm working of the particle dispersion type duplex structure.
  • FIG. 10 shows the relationship between the hardness of the warm working material and the aging temperature.
  • a secondary hardening element such as Mo has been added
  • FIG. 11 shows an example of observation of the ultrafine fiber structure formed in the central part of the plate material with warm rolling at 650°C of N steel by SEM.
  • FIG. 12 shows an example of observation of the ultrafine fiber structure formed in the surface layer part of the locally intense strained rod material by SEM.
  • Table 8 shows the results of the hydrogen embrittlement resistance characteristic test.
  • Results of hydrogen embrittlement resistance characteristics evaluation Steel type Tensile strength (GPa) *1 0.7 TS (GPa ) Diffusive hydrogen content (ppm) Notch strength (GPa) *2 Notch strength (GPa) *3 Hydrogen embrittlement resistance characteristics Ex. 1 A 1.66 1.16 0.36 2.41 1.76 AA Ex. 3 B 1.86 1.30 0.36 2.59 1.78 AA Ex. 7 C 2.08 1.46 0.32 2.17 1.63 AA Comp. Ex. 14 C 1.80 1.26 0.34 2.22 1.79 AA Comp. Ex. 12 C 2.06 1.44 0.25 1.91 0.73 CC Comp. Ex. 16 G 1.79 1.25 0.31 2.00 0.77 CC Comp. Ex.
  • a notch tensile test of a steel charged with hydrogen in an amount of about 0.3 mass ppm was carried out by a slow strain rate tensile test.
  • the hydrogen embrittlement resistance characteristics were evaluated according to whether the notch tensile strength at this step is 0.7 time or more the tensile strength of a smooth tensile test piece not charged with hydrogen.
  • Comparative Example 14 is a steel for a high strength mechanical structure excellent in delayed fracture invented in Japanese Patent Application No. 2001-264399 .
  • a high strength steel multiphased by fine dispersion of a small amount of dispersed second-phase particles is provided.
  • an ultrahigh strength steel which is hard to soften and is hard to form is applied with prescribed deformation in a temperature range in which the deformation resistance is reduced and no cracks occur in the material, to be formed into a prescribed shape (thin plate, thick plate, wire rod, or component).
  • a prescribed shape thin plate, thick plate, wire rod, or component.
  • conventional spheroidizing and quench hardening and tempering treatments after component forming are omitted.
  • an ultrafihe duplex phase structure is developed into a fibrous form.
  • a high strength steel largely improved in the ductility, particularly the toughness, and the hydrogen embrittlement resistance characteristics in the relation of trade-off balance with high strength, and a member thereof.

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Claims (9)

  1. Stahlmaterial mit einer Zusammensetzung, bestehend aus:
    C: mindestens 0,08 Gew.-% und nicht mehr als 0,70 Gew.-%,
    Si: mindestens 0,05 Gew.-% und nicht mehr als 2,5 Gew.-%,
    Mn: mindestens 0,05 Gew.-% und nicht mehr als 3,0 Gew.-%,
    Cr: mindestens 0,01 Gew.-%,
    Al: 0,5 Gew.-% oder weniger,
    O: 0,3 Gew.-% oder weniger,
    N: 0,3 Gew.-% oder weniger, und
    gegebenenfalls einem oder mehreren Elementen, ausgewählt aus:
    Mo: 5,0 Gew.-% oder weniger,
    W: 5,0 Gew.-% oder weniger,
    V: 5,0 Gew.-% oder weniger,
    Ti: 3,0 Gew.-% oder weniger,
    Nb: 1,0 Gew.-% oder weniger,
    Ta: 1,0 Gew.-% oder weniger,
    Ni: 0,05 Gew.-% oder mehr, und
    Cu: 2,0 Gew.-% oder weniger, und
    wobei der Rest im Wesentlichen aus Fe und unvermeidbaren Verunreinigungen besteht;
    wobei das Stahlmaterial ein Teilchendispersionfaserkorngefüge mit einer <011>-Textur parallel zur Laufrichtung (RD) und eine Zugfestigkeit von 1,2 GPa oder mehr aufweist,
    wobei das Teilchendispersionfaserkorngefüge ein Matrixgefüge aufweist, das faserige Ferrit-Kristalle mit einem mittleren Korndurchmesser der Nebenachse von 3 µm oder weniger aufweist,
    wobei das Mikrogefüge des Stahlmaterials Carbidteilchen enthält, die in dem Matrixgefüge des Stahls in einem Volumenanteil von 7 x 10-3 oder mehr und 12 x 10-2 oder weniger dispergiert sind, und der mittlere Durchmesser der Carbidteilchen 0,5 µm oder weniger beträgt.
  2. Stahlmaterial nach Anspruch 1, wobei der mittlere Korndurchmesser von Nebenachsen der faserigen Ferrit-Kristalle 1 µm oder weniger beträgt.
  3. Stahlmaterial nach Anspruch 1 oder Anspruch 2, wobei der mittlere Korndurchmesser von Nebenachsen der faserigen Ferrit-Kristalle 0.5 µm oder weniger beträgt.
  4. Stahlmaterial nach einem der Ansprüche 1 bis 3, wobei der mittlere Korndurchmesser von Hauptachsen der Carbidteilchen 0.1 µm oder weniger beträgt.
  5. Stahlplatte, hergestellt aus dem Stahlmaterial nach einem der Ansprüche 1 bis 4, wobei das Teilchendispersionfaserkorngefüge in mindestens einer Oberflächenschicht der Stahlplatte ausgebildet ist.
  6. Stahlwalzdraht, hergestellt aus dem Stahlmaterial nach einem der Ansprüche 1 bis 4, wobei das Teilchendispersionfaserkorngefüge in mindestens einer Oberflächenschicht des Stahlwalzdrahts ausgebildet ist.
  7. Stahlbolzen, hergestellt aus dem Stahlmaterial nach einem der Ansprüche 1 bis 4, wobei das Teilchendispersionfaserkorngefüge in mindestens einer Oberflächenschicht eines Schraubenabschnitts des Stahlbolzens ausgebildet ist.
  8. Stahlgegenstand, hergestellt durch Schneidbearbeitung des Stahlmaterials nach einem der Ansprüche 1 bis 4, der Stahlplatte nach Anspruch 5, des Stahlwalzdrahts nach Anspruch 6 oder des Stahlbolzens nach Anspruch 7.
  9. Verfahren zur Herstellung eines Stahlmaterials nach einem der Ansprüche 1 bis 4,
    wobei das Verfahren die Schritte umfasst von:
    Tempern eines Stahls unter solchen Bedingungen, dass der durch die folgende Gleichung (1) ausgedrückte Parameter λ : λ = T logt + 20 T ; Temperatur k , t ; Zeit Stdn
    Figure imgb0012
     1,4 x 104 oder mehr bei einer Temperatur im Bereich von 350°C oder mehr und einem Ac1-Punkt oder weniger beträgt und die Vickers-Härte (HV) bei Raumtemperatur gleich oder größer ist als die durch die folgende Gleichung (2) ausgedrückte Härte H: H = 5,2 − 1,2 × 10 − 4 λ × 10 2
    Figure imgb0013
     wobei der Stahl besteht aus:
    C: mindestens 0,08 Gew.-% und nicht mehr als 0,70 Gew.-%,
    Si: mindestens 0,05 Gew.-% und nicht mehr als 2,5 Gew.-%,
    Mn: mindestens 0,05 Gew.-% und nicht mehr als 3,0 Gew.-%
    Cr: mindestens 0,01 Gew.-%,
    Al: 0,5 Gew.-% oder weniger,
    O: 0,3 Gew.-% oder weniger,
    N: 0,3 Gew.-% oder weniger, und
    gegebenenfalls einem oder mehreren Elementen, ausgewählt aus:
    Mo: 5,0 Gew.-% weniger,
    W: 5,0 Gew.-% oder weniger,
    V: 5,0 Gew.-% oder weniger,
    Ti: 3,0 Gew.-% oder weniger,
    Nb: 1,0 Gew.-% weniger,
    Ta: 1,0 Gew.-% oder weniger,
    Ni: 0,05 Gew.-% oder mehr,
    Cu: 2,0 Gew.-% oder weniger,
    wobei der Rest im Wesentlichen aus Fe und unvermeidbaren Verunreinigungen besteht;
    wobei das Mikrogefüge des Stahls 80 %, bezogen auf das Volumen, oder mehr eines Matrixgefüges enthält, das aus einem einheitlichen Gefüge aus Martensit oder Bainit oder einem Mischgefüge aus Martensit oder Bainit besteht,
    wobei in dem Matrixgefüge des Stahls Carbidteilchen in einem Volumenanteil von 7x10-3 oder mehr und 12x10-2 oder weniger dispergiert sind und der mittlere Durchmesser der Zweitphasenteilchen 0.5 µm oder weniger beträgt;
    anschließendes Warmumformen des getemperten Stahls, indem der Stahl einer Dehnung von 0,7 oder mehr bei einer Temperatur im Bereich von 350°C oder mehr und einem Ac1-Punkt von -20°C oder weniger unterzogen wird, und
    gegebenenfalls Altern des warmumgeformten Stahls bei einer Temperatur im Bereich von 350°C oder mehr und einem Ac1-Punkt oder weniger.
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