JP5344454B2 - Steel for warm working, warm working method using the steel, and steel and steel parts obtained thereby - Google Patents

Abstract

There are provided a steel for warm working, to be subjected to warm working as various structures, components of cars, and the like, a warm working method thereof, and a steel material and a steel component obtainable from the warm working method. [Solving Means] A steel is to have a particle dispersion type fiber structure formed in the matrix by warm working. The steel is characterized in that the total amount of the dispersed second-phase particles at room temperature is 7 × 10 -3 or more in terms of volume fraction, and the Vickers hardness (HV) is equal to or larger than the hardness H of the following equation (2): H = (5.2 - 1.2 × 10 -4 ») × 10 2 ... (2) when the steel is subjected to any of annealing, tempering, and aging treatments in the as-unworked state under conditions such that a parameter » expressed by the following equation (1): » = T(logt + 20) (T; temperature (K), t; time (hr)) ... (1) is 1.4 × 10 4 or more in a prescribed temperature range of 350°C or more and Ac1 point or less. This steel is taken as the steel for warm working.

Description

  The present invention relates to steel used by being processed into various structures, automobile parts, and the like. More specifically, the present invention relates to steel for warm working to be used for warm working, its warm working method, and the warm working steel. The present invention relates to steel materials and steel parts obtained by a processing method.

  In recent years, with the increase in the size of structures and the weight reduction of automobile parts, there has been a demand for realization of tougher and higher performance high strength steel than ever before. Conventional measures for improving the toughness of steel include (1) reduction of impurity elements such as P and S that cause embrittlement, (2) refinement and reduction of inclusions, and (3) addition of alloying elements. (4) Reduction of carbon, (5) Crystal grain refinement, (6) Refinement of second phase dispersed particles such as carbide particles are generally known.

  In particular, crystal grain refinement is attracting attention because it has the effect of reducing stress concentration at grain boundaries and diluting impurity elements at grain boundaries, and can increase brittle fracture stress at the same time as yield stress increases. For example, recently, attempts have been made to achieve high strength and long life of steel by reducing the ferrite grain size to 1 μm or less in a low carbon steel considering resource saving and recyclability.

  However, research related to grain refinement of low-carbon ferritic steels so far has concentrated on a strength level of 1000 MPa or less (for example, Non-Patent Document 1-2; Patent Document 1-2). In order to obtain high strength of 1000 MPa or more only by refining ferrite grains, it is necessary to make the grains ultrafine to 0.5 μm or less, and 0.5 μm or less in the thermomechanical processing of steel assuming mass production. This is because it is extremely difficult to miniaturize the material. On the laboratory scale, ultrafine grains of 0.5 μm or less are obtained by super-strong processing methods such as powder metallurgy MM (Non-patent Document 3) and ARB (Non-patent Document 4). Such ultra-fine grain steel shows almost no uniform elongation and non-uniform deformation due to necking becomes the main component of elongation, resulting in a significant decrease in ductility. This same early plastic instability has also been confirmed in pure iron wires strengthened by dislocation by wire drawing (Non-Patent Document 5).

In general, increasing the strength of steel has the problem of significantly reducing various properties such as ductility, toughness, delayed fracture resistance, fatigue characteristics, and formability. In particular, in the low-alloy martensitic steel with high versatility, when strengthened to 1.2 GPa or more, the toughness, delayed fracture resistance and the like are remarkably lowered, so that the practical application of high-strength steel is greatly hindered. Therefore, it is strongly desired to simultaneously achieve high strength and toughness of low alloy steel and improvement of delayed fracture resistance. However, according to the conventional knowledge, as means for improving the fracture characteristics of the low alloy martensitic steel, (a) high temperature tempering avoiding the temper embrittlement temperature range around 500 ° C., (b) old austenite Although refinement of grains, (c) ausfoam, (d) fibrous organization, or a combination of these may be considered, there are the following problems with the application of these means.
(A) High-temperature tempering High-temperature tempering is performed at about 550 ° C. or more and A1 point or less. According to this, (1) the internal stress introduced by quenching can be significantly reduced with the recovery of dislocations. ) There are advantages such as being able to non-align (spheroidize) matched precipitated carbides (for example, film-like cementite) that reduce fracture toughness. For this reason, tempering is usually performed at around 650 ° C. in mechanical structural steels that particularly require toughness. However, in such a temperature range, the second phase dispersed particles easily grow during tempering, so that the strength of the steel is inevitably lowered. Conventionally, a technique has been adopted in which a large amount of carbon is added to increase the precipitation amount of carbides to increase the strength, but the toughness is lowered. Therefore, there is a limit to increasing the strength only by high-temperature tempering. The high strength can be achieved even by high temperature tempering, limited to steels added with a large amount of special alloy elements such as maraging steel (Non-Patent Document 6-10).
(B) Crystal grain refinement It is indispensable to refine the prior austenite grains so that sufficient toughness can be ensured when steel is strengthened. As a method for refining austenite grains, there are (1) a method by recrystallization of processed austenite and (2) a method using phase transformation. Among them, the work heat treatment in which the martensite structure, which is classified as the latter, is processed in the cold or warm region and then austenitized, can most effectively refine the austenite grains (Non-patent Documents 11 and 12). It was thought. For example, it is known that the toughness of tempered martensitic steel is improved (Non-patent Document 13) and delayed fracture characteristics are improved (Patent Document 4) by refining austenite grains of several μm or less (Patent Document 4). 3). In the refinement of austenite, the grain growth rate increases as the crystal grains become finer, so how to suppress the grain growth during austenite was a particularly important point. Therefore, conventionally, dispersion of pinning particles effective for suppressing the growth of austenite, reduction of austenitizing temperature, rapid austenitizing using high-frequency heating, and the like have been generally applied. However, it is extremely difficult to suppress the growth of ultrafine austenite grains, and in actuality, the grain refinement has reached its peak at about several μm. Further, if the crystal grains are made too fine, there is a problem that diffusion phase transformation at the grain boundary is promoted and it becomes difficult to burn, and the process window for austenite grain ultrafine refinement is relatively narrow.
(C) Ausfoam Ausfoam is a process in which austenitized steel is rapidly cooled to the metastable austenite region, processed at that temperature, quenched and martensitic or bainite transformed, and then tempered. Can be strengthened without significantly impairing its toughness. In this ausform, there are effects such as (1) miniaturization of packets and blocks that are effective crystal grains, (2) dislocation takeover from processed austenite to martensite, and (3) dislocation pinning by carbon atoms or carbides. It is thought that the steel is strengthened by overlapping. Recently, improved ausforms that process in the high temperature metastable austenite region have been applied to medium carbon low alloy steels, and improvements in fatigue and delayed fracture properties have been reported. The main factors for improving the characteristics by the improved ausfoam are the refinement of the base structure, the suppression of the formation of coarse grain boundary cementite by the introduction of grain boundary irregularities (Non-Patent Document 14) and the formation of texture (Non-Patent Document 15). Is considered. However, since austenite is an austenitic structure process, it was necessary to strictly adjust the alloy components and thermomechanical conditions so that the metastable austenite phase does not cause pro-eutectoid ferrite transformation or pearlite transformation during processing. . In addition, since there is a problem of causing burning cracks during cooling after processing, the applied members are limited to simple shapes such as plates and bars.
(D) Fibrous organization It is also effective to generate a fiber structure inside by cold or warm processing for strengthening steel. This is already the case with ausformed steel (Non-Patent Documents 16 and 17), high-strength low-carbon wire for cold wire drawing (Patent Document 5), piano wire, pure iron wire (Non-Patent Document 5), etc. Proposed.

  With regard to the processing of steel materials, cold-working has become the main process for forming parts today because it can mass-produce parts with complex shapes such as bolts with dimensional accuracy. However, for steel materials having a tensile strength exceeding 1.2 GPa, cold forging is extremely difficult due to the strength, and thus, for example, a member in which a fiber structure is formed by the cold forming process as described above. Was limited to wire.

  On the other hand, many attempts have been made for warm working in a two-phase region of a ferrite phase and a carbide below the Ac1 point. For example, a high-strength member and a high-strength steel material are prepared, and the material is warm-processed so as to produce a member having a desired geometric shape in a state in which the strength characteristics of the material are substantially retained or enhanced. There has been known a method and a forming method for forming a structure and producing a high-strength steel structural member having at least a tensile strength of 1 GPa (Patent Document 6).

  Furthermore, a material having an ultrafine structure is warm-worked or cold-worked, a steel material made of expanded ferrite grains having a minor axis of 3 μm or less is used as the material, and only the forming is performed without any tempering treatment. There is also known a method for producing a molded product characterized by not performing quality treatment (Patent Document 3).

  Further, in steels having a multiphase structure such as a tempered martensite structure, warm working has been applied for the purpose of obtaining a processed structure before quenching treatment for refining reverse transformed austenite (non- Patent Documents 11 and 12). Since the strength of the steel is achieved by tempering the steel subsequent to the warm working, no attempt has been made to use the material in the warm working of the tempered martensite structure.

  Furthermore, in warm straightening of high carbon steel having a carbon content of 0.7 mass% or more, a wire rod exceeding 1.8 GPa is obtained, but the elongation of the wire rod is as low as about 6% (Non-patent Document 18). ).

  As described above, prior austenite grain refinement and ausforming are important toughening techniques for steel, and their research and inventions are enormous. However, in these processes, quenching and tempering are fundamental, and the increase in strength is limited by the problems of hardenability and tempering cracks and the problem of temper embrittlement. Further, as the strength is increased, the amount of second phase dispersed particles such as carbide necessary for strengthening also increases, so that softening by spheroidizing annealing or the like is difficult. In addition, particularly when the carbide is coarsened by annealing, there is a problem that cracks are generated in the material during the process of forming the material into parts by cold forging or the like. Therefore, as long as the conventional high-strength process by quenching and tempering is used, the tensile strength is 1.2 GPa or more, and the characteristics of high-strength steel of 1.5 GPa or higher that makes it difficult to soften are greatly improved. It was impossible to make it practical.

  In addition, research and inventions related to warm working so far are mainly aimed at forming into a member and preparing a pre-structure. Therefore, starting with a relatively soft base structure such as ferrite, pearlite structure with low deformation resistance, or martensite structure tempered at high temperature, warm processing is performed under conditions where deformation resistance is low. Most of them were applied. In addition, the fine multiphase structure is not taken into consideration in consideration of the dispersion state and thermal stability of the second phase dispersed particles, and the tensile strength is 1.2 GPa or more after warm working, and the ductility and toughness. A high strength member excellent in delayed fracture characteristics has not been realized. In particular, for multiphase steels such as tempered martensite steel having a tensile strength of 1.2 GPa or more at room temperature, there is a risk that warm working cannot be performed due to its strength. It seemed impossible.

  Therefore, the present invention solves the problems as described above, warm-works high-strength steel having a tensile strength of 1.2 GPa or more, excellent ductility and delayed fracture resistance, and dramatically improved toughness. It aims at providing the steel for warm work which can produce | generate the particle dispersion type | mold fibrous structure | tissue for obtaining by this, and the warm work method using the same. Another object of the present invention is to provide a steel material such as a steel plate and a steel bar, and a steel part such as a bolt and a machined part provided with the above characteristics.

  The inventor has intensively studied to solve the above problems, and as a result, has made the following invention.

First: the chemical composition is
C: 0.70 mass% or less,
Si: 0.05 mass% or more and 2.5 mass% or less,
Mn: 0.05 mass% or more and 3.0 mass% or less,
Cr: 0.01 mass% or more and 2.01 mass% or less,
Al: 0.5 mass% or less,
O: 0.3 mass% or less,
N: 0.3 mass% or less,
Mo: 0.002 mass% or more and 5.0 mass% or less,
The balance a steel material which is substantially Fe and incidental impurities, wherein the steel consists of grain disperse fibrous grain structure exhibiting a <011> // RD (rolling direction) texture, fibers constituting the base structure Vickers of the steel material at room temperature is obtained by finely dispersing the second phase particles in the matrix structure at a volume ratio of 7 × 10 ⁻³ or more and 12 × 10 ⁻² or less with an average particle size of minor axis of the ferrite ferrite crystal of 3 μm or less. A steel material having a hardness of HV 3.7 × 10 ² or more.

  Second: The steel structure is characterized in that the base structure is composed of a fibrous ferrite crystal having a minor axis average particle diameter of 1 μm or less.

  Third: The steel structure is characterized in that the base structure is composed of a fibrous ferrite crystal having a minor axis average particle size of 0.5 μm or less.

Fourth: A steel material wherein the average particle size of the major axis of the second phase dispersed particles is 0.1 μm or less.
Fifth: The above steel material,
W: 5.0 mass% or less,
V: 5.0 mass% or less,
Ti: 3.0 mass% or less,
Nb: 1.0 mass% or less,
Ta: 1.0 mass% or less
A steel material further comprising one or more selected from the group consisting of:
Sixth: The above steel material,
Ni: 0.05 mass% or more and 9 mass% or less,
Cu: 2.0 mass% or less
A steel material characterized by further containing one or two of the above.

Seventh : A steel for warm working for creating the above steel material by warm working, and a parameter λ represented by the following formula (1) is 1. in a predetermined temperature range of 350 ° C. or higher and Ac1 point or lower. Second phase particles are generated by performing heat treatment of any one of annealing , tempering, or aging treatment under the condition of 4 × 10 ⁴ or more, and the volume ratio of the second phase particles at room temperature after this heat treatment is in 7 × ^{10 -3} or more 12 × ^{10 -2} or less, and steel Vickers hardness (HV) is the following formula (2) in hardness H or more warm working steel, characterized in that the. λ = T (logt + 20) (T; temperature (K), t; time (hr)) (1)
H = (5.2-1.2 × 10 ⁻⁴ λ) × 10 ² (2)
Eighth : A steel for warm working characterized in that 80% by volume or more of the base structure is a single structure of martensite or bainite, or a mixed structure thereof.

9: a steel plate obtained by processing between the temperature of the above warm working steel, steel, characterized in that the particle dispersion type fibrous crystal structure at least on its surface layer portion is generated Board material.

Tenth: A steel bar obtained by warm-working the above-mentioned warm-working steel, wherein a particle-dispersed fibrous crystal structure is generated at least on the surface layer thereof. Bar material.

11th: A steel bolt obtained by warm-working the above-mentioned steel for warm-working, wherein the steel-coated bolt has a particle-dispersed fibrous crystal structure formed at least on its surface layer .

12: the steel, the steel plate, a method for producing a steel bar or steel bolts, the following formula in a predetermined temperature range below Ac1 point 350 ° C. or higher for the warm working steel (1) A warm working method characterized in that a desired shape is obtained by warm working with a strain of 0.7 or more under the condition that the parameter λ expressed by is 1.4 × 10 ⁴ or more.
λ = T (logt + 20) (T; temperature (K), t; time (hr)) (1)
13th : A steel processed product obtained by cutting a steel material, wherein the steel material is any one of the above steel material, a steel plate material, a steel bar material, or a steel bolt.

According to the first invention, a steel material having not only high toughness but also improved secondary workability by the formation of the fine fiber structure is realized. In addition, according to the first invention, it is possible to achieve high strength of steel obtained when subjected to warm working by an alloy composition excellent in economic efficiency and recyclability.

  According to the second invention, a dense fiber structure having an average short axis average interval of 1 μm or less and according to the third invention an average short axis interval of 0.5 μm or less is developed. A steel material with higher toughness and workability is realized.

According to the fourth invention, by controlling the average particle diameter of the long axis of the second phase dispersed particles to be 0.1 μm or less, further enhancement of strength and toughness can be achieved by dispersing a small amount of the second phase dispersed particles. realizable.
According to the fifth aspect of the invention, it is possible to disperse the second phase dispersed particles that are finer and excellent in hydrogen trapping properties, and to increase the strength of the steel material obtained when subjected to warm working and in a low temperature range. Toughness and delayed fracture resistance can be greatly improved.
According to the sixth invention, the toughness can be further improved to a low temperature range.

According to the seventh and eighth inventions, by controlling the softening resistance at the time of warm working and the volume fraction of the second phase dispersed particles, when subjected to warm working, the tensile strength is dramatically improved. A particle-dispersed fibrous structure was produced. As a result, a steel for warm working is realized in which the steel material after warm working retains a tensile strength of 1.2 GPa or more at room temperature while dramatically improving its toughness.

  In addition, by using a martensitic transformation or bainite transformation to form an ultra-fine multiphase structure in which second-phase dispersed particles such as carbide particles are finely dispersed, it is efficiently incorporated inside when subjected to warm working. It becomes possible to generate the fiber structure. At the same time, the delayed fracture resistance can be greatly improved.

According to the ninth and tenth inventions, not only has high toughness and tensile strength, but also has secondary workability, so that it can be used for manufacturing various parts and products, and the practicality is dramatically improved. Steel plates and bar wires have been realized.

According to the first aspect of the invention, a bolt excellent in impact resistance and delayed fracture resistance in which a fiber structure is generated at the bottom of the threaded portion where stress is concentrated is realized.

According to the twelfth invention, a high-toughness can be obtained by generating a fiber structure while processing the warm-working steel into a desired shape. In addition, since the warm processing equipment conventionally utilized as an equipment can be utilized, it has extremely high practicality.

According to the thirteenth invention, even a high-strength part having a complicated shape is provided as being excellent in impact resistance and delayed fracture resistance.

  As described above, the present invention is applicable to high-strength steel that has been made into a multi-phase by fine dispersion of a small amount of second-phase dispersed particles, particularly to ultra-high-strength steel that is difficult to soften and difficult to form. After the conventional spheroidizing annealing and component molding, the material is molded into a predetermined shape (thin plate, thick plate, bar wire, part) by giving a predetermined deformation in a temperature range where the material does not crack and does not crack in the material High-strength steels that significantly improve the ductility, especially toughness and delayed fracture resistance, which have a trade-off balance between high strength by developing an ultrafine multiphase structure into a fibrous shape while omitting quenching and tempering Providing a member.

The said effect is based on the following mechanisms.
(A) Ultrafine grain refinement and formation of a fibrous base structure by warm working If the material satisfies certain conditions, it has far greater toughness and delayed fracture resistance than conventional ausfoam steel. The knowledge that an excellent particle-dispersed fiber structure can be formed on a member was obtained. That is, the pinning effect by the fine dispersion or precipitation of the second phase dispersed particles is effectively utilized, and the recovery of dislocations introduced by deformation occurs moderately, but the material does not undergo primary recrystallization or significant grain growth. When the crystal grains are refined by applying a predetermined strain by deforming, an ultrafine grained multiphase structure having low internal stress and no crack initiation point can be created. In particular, in such ultrafine grains, by developing a fiber structure with a narrower crystal grain boundary interval, it is possible to significantly increase fracture toughness by suppressing not only the occurrence of cracks but also the propagation of cracks.
(B) Refinement of coarse second phase Even coarse second phase dispersed particles that cause cracking in cold working can be relatively easily deformed without cracking in warm working. Therefore, by utilizing the decomposition and reprecipitation of the second phase dispersed particles generated during processing, not only spheroidizing the coarse film-like precipitate considered to be the cause of intergranular cracking but also finely dispersing it. Can be used for strengthening.
(C) Ultrafine dispersion of alloy carbides and intermetallic compounds, etc. Alloy elements with high carbide forming ability such as Mo, V, W, Ta, Ti, Nb, etc. are unique in Mo ₂ C. , V ₄ C ₃ , W ₂ C, TaC, NbC, TiC and other nano-sized alloy carbides are formed in the temperature range of 500 ° C. to 600 ° C. Therefore, the addition of these alloying elements can increase the strength of steel It is effective for conversion. The maximum value of precipitation strengthening by these nano-sized alloy carbides is obtained in the transition region where the strengthening mechanism changes from Cutting to the Owanan mechanism. Toughness decreases. For this reason, steel is usually tempered to a sufficiently overaged state of these carbides, even at the expense of some strength of the steel. On the other hand, if dynamic precipitation of these alloy carbides by warm working is utilized, it is possible to cause inconsistent precipitation of carbides without much carbide growth even in the above-mentioned precipitation transition temperature range. In other words, precipitation strengthening of alloy carbide by the Orowan mechanism can be used to the maximum extent possible. Similar effects can also be expected for precipitation of intermetallic compounds, nitrides, oxides, Cu particles, and the like composed of the above alloy elements and Ni, Al, and the like.

  The present invention has the features as described above. The requirements of the present invention will be described in detail below.

The steel for warm working of the present invention is annealed without being processed under the condition that the parameter λ represented by the following formula (1) is 1.4 × 10 ⁴ or more in a predetermined temperature range of 350 ° C. or more and Ac1 point or less. And an alloy element that generates second phase particles by performing heat treatment of either one of tempering or aging treatment, and the volume ratio of the second phase particles at room temperature after this heat treatment is 7 × 10 ^−3. It is characterized in that it is 12 × 10 ⁻² or less and its Vickers hardness (HV) is not less than the hardness H of the following formula (2).
λ = T (logt + 20) (T; temperature (K), t; time (hr)) (1)
H = (5.2-1.2 × 10 ⁻⁴ λ) × 10 ² (2)
Thus, since the steel for warm working of the present invention changes the dispersion state of the second phase dispersed particles and the base structure during the warm working applied thereto, it is a tempering process that simulates the heat history of warm working. It is comprised by setting the minimum of Formula (2) with respect to the hardness (structure | tissue) of the unprocessed material obtained. That is, as described below, the tissue state is represented by hardness.

(A) Structure of steel for warm working In order to achieve high strength and toughening of the duplex structure steel simultaneously by warm working, strengthening by dispersing the second phase dispersed particles as small as possible and the base It is important that the structure can be refined and the fiber textured simultaneously. In order to achieve this ultrafine multiphase structure, the fine dispersion or fine dispersion ability of the second phase dispersed particles in the steel for warm working as the material is important.

In the present invention, regarding the fine dispersion or fine dispersion ability of the second phase dispersed particles,
(I) The tempering treatment is already performed in the steel for warm working and the second phase dispersed particles are dispersed. (Ii) The second phase dispersed particles are not dispersed in the steel for warm working. One type or two or more types of second phase dispersed particles are precipitated therein, and a particle dispersed fiber structure is formed after processing (iii). Are dispersed, but three types of precipitation of other particles during warm working can be considered.

  The dispersion (precipitation) strengthening by the second phase dispersed particles depends on the dispersion state such as the volume ratio of the second phase dispersed particles, the size of the particles, the hardness and the shape. When dispersion strengthening is based on the Orowan mechanism, the particle diameter (d) is smaller than the following formula (A) (“Steel Precipitation Control Metallurgy Frontline (Japan Iron and Steel Institute) (2001) P.69”), and the volume fraction ( The dispersion strengthening amount increases as f) increases. That is, the dispersion state (and dispersibility) of the second phase dispersed particles has a close relationship with the hardness.

Δσ = (3.2 Gb) / [(0.9f ^−1/2 −0.8) d] (A)
Here, G is a steel rigidity of 80 GPa, and b is a Burger spectrum of 0.25 nm.
However, if the particle becomes too smaller than a certain critical particle size, the dislocation is not pinned by the particle, and the particle is sheared by the dislocation, so that the Orowan mechanism is not established. In the so-called Cutting mechanism in which particles are sheared by dislocation, the amount of dispersion strengthening increases as the particle diameter increases. That is, the maximum dispersion strengthening amount can be obtained with the minimum particle diameter at which the Orowan mechanism is established. The minimum particle size at which the maximum dispersion strengthening can be achieved depends on the hardness of the particles and decreases in inverse proportion to the hardness of the particles (frontier of precipitation control metallurgy of steel (Japan Iron and Steel Institute) (2001) p. 69) . Therefore, when compared at the same volume ratio, the harder particles have a smaller minimum particle diameter at which the Owanan mechanism is established, and therefore the maximum amount of particle dispersion strengthening is increased.

For example, it is known that TiC has a high hardness among alloy carbides and has a small density, so that effective dispersion particle strengthening can be performed. Assuming that 7 nm is obtained as the minimum particle size applicable to the Orowan mechanism with TiC, about 0.9 GPa (TS (GPa) ≈0.0032 HV, HV 2.8 ×) with a volume fraction dispersion of 7 × 10 ^−3. A particle dispersion strengthening amount of 10 ² ) can be expected. Incidentally, when the density of TiC is 4.94 Mg / m ³ , the atomic weight of Ti is 47.9, and the atomic weight of C is 12, the Ti necessary for precipitating TiC with a volume ratio of 7 × 10 ⁻³ is 0.35 mass%, C Is 0.087 mass%. In addition, since the strength of the base of the practical ferritic steel is about 0.3 GPa (about HV 0.9 × 10 ² ), the room temperature strength of the steel in which the TiC is dispersed in the ferrite base is 1.2 GPa or more (HV 3.7). X10 ² or more). Therefore, considering the ideal dispersion state for TiC, the dispersion particles to which the Orowan mechanism can be applied have an HV of 3.7 × 10 ² only by dispersion strengthening at a small volume ratio of 7 × 10 ⁻³ if the size is 7 nm. You will be fully satisfied. The same effect can be expected for the second phase dispersed particles composed of carbonitride, intermetallic compound, oxide, Cu particles and the like. As such second phase dispersed particles, specifically, for example, carbides such as Mo ₂ C, V ₄ C ₃ , W ₂ C, TaC, NbC, TiC, Fe ₃ O ₄ , Fe ₂ O ₃ , Al ₂ O ₃ , Cr ₂ O ₃ , SiO ₂ , Ti ₂ O _{3 and} other oxides, AlN, CrN, TiN and other nitrides, Ni ₃ Ti, NiAl, TiB, Fe ₂ Mo, Ni ₃ Nb, Ni ₃ Intermetallic compounds such as Mo, metal particles such as Cu particles, and the like can be considered.

It is known that metal carbide particles such as Mo and Ti are generally about 10 nm in size, and that high strength can be effectively achieved even with a small amount of dispersion having a volume ratio of less than 10 × 10 ⁻³ . However, the size of the second phase dispersed particles and the distribution in the matrix structure also vary due to segregation of alloy elements and the like. Therefore, in the present invention, even if there is variation in the distribution of the second phase dispersed particles, the volume fraction of the second phase dispersed particles at room temperature is considered so that a fine crystal structure can be stably obtained by warm working. ^Is defined as 7 × 10 ⁻³ or more. For low alloy martensitic steels and baitnite steels, the second phase dispersed particles are taken into account when the average particle size of general cementite (Fe ₃ C) before warm working is several tens of nanometers or more. The volume ratio is preferably 20 × 10 ⁻³ or more.

In addition, when the strength is increased, the upper limit of the volume ratio of the second phase dispersed particles is set to 12 × 10 ⁻² or less in consideration of toughness. Further, the particle dispersion strengthening by the Orowan mechanism is expected to become remarkable in the region of several tens of nm or less from the formula (A), and 1 in the dispersed state of the second phase dispersed particles having an average particle diameter larger than 0.5 μm. It is difficult to obtain a strength of 2 GPa or more. Therefore, it is desired for the steel for warm working that the average particle size of the second phase dispersed particles is 0.5 μm or less, more preferably 0.1 μm or less.

However, the above conditions are based on the premise that the second phase dispersed particles do not grow even in a temperature range of 350 ° C. or higher in the third stage or higher. In other words, in order to have a strength of 1.2 GPa or more even after warm processing, it is necessary that the second phase dispersed particles notably decrease in strength due to remarkable Ostwald growth in addition to the matrix structure during and after heating, processing. It becomes a condition. Therefore, when the thermal stability of the tissue is evaluated using λ represented by the following equation (1), which is generally known as a tempering parameter, as an index, λ ≧ 1 in a predetermined temperature range of 350 ° C. or more and Ac1 point or less. Temper softening resistance such that the Vickers hardness (HV) at room temperature when tempering is performed without processing under the condition of 4 × 10 ⁴ is equal to or higher than the hardness H given by the following formula (2) It can be considered that the pre-processed structure, that is, the necessary and sufficient condition for the warm-working steel of the present invention is shown.

λ = T (logt + 20) (1)
Here, T is temperature (K), and t is time (h).

H = (5.2-1.2 × 10 ⁻⁴ λ) × 10 ² (2)
The term “in the predetermined temperature range” means that the above condition should be satisfied at any temperature from 350 ° C. to Ac1 point, and means that the above condition need not be satisfied over the entire temperature range. That is, in the case of aging or tempering treatment, when the material undergoes remarkable age hardening or secondary hardening and becomes a hardness H or more only in a certain temperature range within the above range, Can be steel.

Here, for example, the grain growth suppressing effect of the TiC will be considered. From the well-known Zener relation (D = B × d / f), TiC is dispersed (d = 7 × 10 ⁻³ μm, volume ratio f = 7 × 10 ⁻³ , B = 4 / 9˜ 4/3) The stable crystal grain size D obtained by the steady grain growth of the ferrite structure is estimated to be about 0.4 to 1.3 μm. That is, since such stable crystal grains are established for normal grain growth of recrystallized grains, in the warm working in a region lower than the recrystallization temperature, the refinement of the base structure by applying a predetermined strain and TiC It is sufficiently expected that a fibrous structure having an average particle diameter of 3 μm or less can be obtained by the two effects of grain boundary pinning.

Thus, in order to obtain an ultrafine multiphase structure having a tensile strength of 1.2 GPa or more by warm working on the basis of precipitation strengthening by the ideal dispersion state of TiC carbide, the volume fraction of the second phase dispersed particles is The hardness of the steel after tempering is HV ≧ (5.2-1.2 × 10 ⁻⁴ λ) under the condition that the lower limit is 7 × 10 ⁻³ and T (logt + 20) ≧ 1.4 × 10 ^4. Having x10 ² is a necessary and sufficient condition for the pre-processed structure. That is, as the steel for warm working, the second phase dispersed particles are finely dispersed or precipitated as the particle dispersion strengthened particles in the base structure, and the structure control for improving the thermal stability of the second phase dispersed particles It is a feature of the invention.

  About the structure | tissue of the steel for warm processing of this invention as mentioned above, since the dispersion state and base structure | tissue of a 2nd phase dispersion particle are variously changed during the process of warm processing, it is limited by the structure | tissue form of room temperature. However, in reality, all steels having a strength of 1.2 GPa or more, excluding steels having a pearlite structure as the main structure, can be considered as warm work steels. As such a material, specifically, for martensite steel (tempered martensite structure), a low alloy steel of JIS-G4053, a spring steel of JIS-G-4801, and a strength level of 2 or higher. Secondary hardened steel, maraging steel, TRIP steel, ausformed steel and the like.

  And the steel for warm work of this invention makes it 80 volume% or more of a base structure to be a single structure of martensite and bainite, or these mixed structures. This is because, in a medium carbon low alloy steel, the width of a block that is considered to be an effective crystal grain of martensite is 1 μm or less (Script Mater., 49 (2003), P. 1157). In addition to being able to efficiently form a fiber structure by performing warm working on a tempered martensite structure in which carbides are finely dispersed, the bainite structure is also a needle-like or plate-like structure form in which carbides are finely dispersed. This is because a fibrous structure can be obtained in the same manner when this is used as a pre-processed structure. In the steel for warm working of the present invention, it is more preferable that the single structure of such martensite and bainite or the mixed structure thereof is 90% by volume or more of the base structure.

In particular, in order to stably maintain a strength of 1.2 GPa or more after warm working, a martensite or bainite structure having a temper softening resistance equal to or higher than that of tempered martensite steel of JIS-SCM430 steel is 80. % Or more is desirable. In addition, 20 volume% or less other than martensite or bainite and mixed structures thereof may be any structure such as ferrite, pearlite, and austenite structure. This is because such a ferrite, pearlite, austenite structure, etc. decomposes / disappears during the warm working heat treatment or changes to a fine structure, so it is judged that there is no problem if it is 20 volume% or less. is there.
(B) Chemical composition As a preferable chemical composition of the steel for warm working of the present invention, C: 0.70 mass% or less, Si: 0.05 mass% or more, Mn: 0.05 mass% or more, Cr: 0.01 mass% As described above, Al: 0.5 mass% or less, O: 0.3 mass% or less, N: 0.3 mass% or less are contained, and the balance is substantially Fe and inevitable impurities. Furthermore, Mo: 5.0 mass% or less, W: 5.0 mass% or less, V: 5.0 mass% or less, Ti: 3.0 mass% or less, Nb: 1.0 mass% or less, Ta: 1.0 mass% or less It can be considered to contain one or two or more selected from the group consisting of Ni, 0.05% by mass or more and Cu: 2.0% by mass or less. Below, the preferable each component structure | tissue of the steel for warm work of this invention is described.

  C: C forms carbide particles and is the most effective component for increasing the strength. However, if it exceeds 0.70 mass%, the toughness is deteriorated, so the content is set to 0.70 mass% or less. In order to sufficiently expect an increase in strength, it is preferable to contain 0.08 mass% or more, more preferably 0.15 mass% or more.

  Si: Si is an element effective for deoxidizing and dissolving in ferrite to increase the strength of steel and finely disperse cementite. Therefore, the content of 0.05% by mass or more including those added as a deoxidizer and remaining in the steel. Considering the workability of the steel material, it is set to 2.5 mass% or less.

  Mn: Mn is an element that is effective for lowering the austenitizing temperature and making the austenite finer, and is effective for suppressing the coarsening of the cementite by hardening into the hardenability and cementite. If the amount is less than 0.05 mass%, a desired effect cannot be obtained. More preferably, 0.2 mass% or more is contained. In consideration of the toughness of the obtained steel material, it is set to 3.0 mass% or less.

  Cr: Cr is an element effective for improving hardenability and is an element having a strong effect of delaying the growth of cementite by being dissolved in cementite. In addition, it is also one of the important elements in the present invention that, by adding a relatively large amount, forms a high Cr carbide that is more thermally stable than cementite and improves corrosion resistance. Therefore, it is necessary to contain at least 0.01 mass%. Preferably it is 0.1 mass% or more, More preferably, 0.8 mass% or more is contained.

  Al: Al is an element effective for deoxidizing and forming an intermetallic compound with an element such as Ni to increase the strength of the steel. However, since excessive addition reduces toughness, it was made into 0.5 mass% or less. In the case where an intermetallic compound of Al and another element, Al nitride, oxide, or the like is not used as the second phase dispersed particles, it is 0.02 mass% or less, and more specifically 0.01 mass% or less. It is preferable.

  If O: O (oxygen) can be finely and uniformly dispersed as an oxide, it effectively acts not as inclusions but as grain growth suppression and dispersion strengthening particles. However, if excessively contained, the toughness is lowered, so the content was made 0.3 mass% or less. When the oxide is not used as the second phase dispersed particles, the content is preferably 0.01% by mass or less.

  If N: N (nitrogen) can be finely and uniformly dispersed as a nitride, it effectively acts as a grain growth inhibiting particle or dispersion strengthening particle. However, if excessively contained, the toughness is lowered, so the content was made 0.3 mass% or less. When nitride is not used as the second phase dispersed particles, the content is preferably 0.01 mass% or less.

  Mo: Mo is an element effective for increasing the strength of steel in the present invention, and not only improves the hardenability of the steel, but also solidifies a small amount in cementite to make the cementite thermally stable. In particular, the steel is strengthened by secondary hardening by completely nucleating alloy carbide on the dislocations in the matrix phase completely separate from cementite. Moreover, the formed alloy carbide is effective for atomization and also for hydrogen replacement. Therefore, preferably 0.1 mass% or more, more preferably 0.5 mass% or more is contained, but it is an expensive element and excessive addition forms coarse undissolved carbides or intermetallic compounds to deteriorate toughness. Therefore, the upper limit of the addition amount is set to 5 mass%. From an economical viewpoint, it is preferable to set it as 2 mass% or less.

  W, V, Ti, Nb and Ta also showed the same effect as Mo, and the upper limit addition amount was determined for each. Furthermore, the combined addition of these elements is effective in finely dispersing the dispersion strengthening particles.

  Ni: Ni is an element that is effective for improving hardenability and is effective for reducing the austenitizing temperature and reducing the austenite, improving toughness, and improving corrosion resistance. Moreover, if it contains an appropriate amount, it is an element that is effective in forming an intermetallic compound with Ti or Al to strengthen precipitation of steel. If the amount is less than 0.01 mass%, the desired effect cannot be obtained. More preferably, 0.2 mass% or more is contained. Although there is no restriction | limiting in particular about an upper limit, Since it is an expensive element, it is preferable to set it as 9 mass% or less.

  Cu: Cu is a harmful element that causes hot brittleness, but if an appropriate amount is added, it causes precipitation of fine Cu particles at 500 ° C. to 600 ° C. and strengthens the steel. When added in a large amount, hot brittleness is caused, so the amount was set to 2 mass% or less, which is almost the maximum solid solution amount in ferrite.

  P (phosphorus) and S (sulfur) are not particularly defined, but P and S are elements that should be removed as much as possible in order to reduce the grain boundary strength, and are each preferably 0.03 mass% or less.

  In addition, it is permissible for elements other than those described above to be contained in various elements as long as the effects of the present invention are not reduced.

(C) Preparation of steel for warm working The method for producing the steel for warm working as described above is, for example, a variety of methods according to the manufacturing method of martensite structure and bainite structure of JIS standard. Can be considered.

(D) Warm working The warm working method of the present invention is a temperature that gives a strain of 0.7 or more to any one of the above-mentioned steels for warm working in a temperature range of 350 ° C or higher and Ac1 point -20 ° C or lower. It is characterized by inter-processing. It is also considered to perform an aging treatment in a temperature range of 350 ° C. or higher and Ac1 point or lower after the warm working. According to such warm processing,
(1) Dislocation recovery occurs moderately, crystal grains can be refined and internal stress can be reduced (2) Diffusion of alloy elements becomes relatively easy, and decomposition and reprecipitation of second phase dispersed particles such as carbides Remarkably occurs and the structure can be refined. (3) The deformation resistance (high temperature hardness) of the steel is remarkably lowered, and an advantage that it can be formed without occurrence of cracks can be obtained.

  More specifically, with respect to such a processing temperature, for example, in the case of a medium carbon low alloy steel used as a general machine structural steel and based on a martensite structure, a tempering third stage in which cementite precipitates. It is possible to set the temperature to 350 ° C. or higher that substantially corresponds to In particular, in order to effectively use alloy carbides, intermetallic compounds, Cu, and the like as the second phase dispersed particles, it is necessary to process in a temperature range of 500 ° C. to 650 ° C., which is the precipitation temperature of these second phase dispersed particles. desirable.

On the other hand, in the portion that has undergone austenite transformation during processing, phase transformation such as pearlite transformation or martensite transformation occurs in the cooling process, and as a result, there is a high possibility that a non-uniform structure that causes cracking is formed. Moreover, the upper limit temperature of processing was made into Ac1 point-20 degreeC also considering the temperature rise by process heat_generation | fever. However, as a combination of the processing temperature and time of the material, when the hardness is arranged by the tempering parameter λ, the Vickers hardness at room temperature is HV 3.7 × 10 ² or less when the material is tempered without being processed. The combination which does not become is preferable in order to obtain the strength of 1.2 GPa or more after warm working. Particularly in processing in a high temperature range, it is necessary to shorten the time required for processing in consideration of the softening resistance of the material and the heating time.

  The degree of tissue development depends on the pre-processed structure, processing temperature and strain. In other words, the amount of strain required varies depending on the pre-processed structure and processing temperature, so the amount of strain cannot be strictly defined here, but when trying to form a fibrous structure inside the material, 0.7 or more, more It is preferable to apply one or more strains. Less than 1 strain amount for warm work steels with martensite and bainite structure in which prior austenite grains are elongated into fine fibers by processing in advance in the non-recrystallization temperature range of austenite A fine fiber structure can be uniformly generated by the application of. However, in an approximate case, the amount of strain is preferably 1 or more, more preferably 1.5 or more.

  At this time, the strain to be applied is not limited to a single process, and may be introduced in a plurality of processes. Further, the processing direction is not always limited to the same direction. Furthermore, the time between passes is not particularly limited. Furthermore, it includes applying a predetermined strain to a specific region (for example, a surface layer that requires high strength, an R portion of a part, or the like) instead of the entire area of the workpiece. However, the actual amount of strain depends on the material properties of the workpiece, the conditions of the roll (mold if it is forged) and the workpiece (such as the type and presence of lubricant), and the roll (the mold if forged). ), The rolling (forging) speed, the rolling (forging) temperature, etc., can only be understood. In particular, when parts are formed by forging, it is essential that non-uniform strain is introduced. Therefore, it is desirable to predict the amount of strain by a highly accurate numerical analysis technique, but generally the rolling reduction is 45% or more in the case of sheet rolling on the premise of a plane strain state, and the cumulative surface reduction in the case of bar rolling. If the rate is 45% or more, it is considered that a strain of 0.7 or more is introduced in the entire area of the workpiece. If the cumulative rolling reduction or cumulative area reduction is 58% or more, it is considered that a strain of 1 or more is introduced throughout the workpiece. However, for example, even if the rolling reduction (area reduction) is less than 45%, a strain of 0.7 or more may be introduced into the entire area of the workpiece or a specific region due to the influence of friction, in that case. It is necessary to quantitatively study the amount of strain introduced by numerical analysis.

(E) Steel material The steel material of the present invention is a steel material obtained by warm-working the steel for warm working as described above, and has a base structure composed of fibrous crystals having a minor axis average particle size of 3 μm or less. And the second phase dispersed particles are finely dispersed in the matrix structure at a volume ratio of 7 × 10 ⁻³ or more at room temperature, and the Vickers hardness at room temperature is HV3.7 × 10 ² or more. . The base structure in the steel material of the present invention is composed of fibrous ferrite crystals having an extensibility (aspect ratio) exceeding 2 and typically an aspect ratio of 5 or more, and the second phase dispersed particles are finely dispersed therein. Can be understood.

  The effect of grain refinement on the mechanical properties of steel is known to be noticeable in the grain region of several μm or less. In the present invention, the average interval between the base structures composed of fibrous crystals (that is, the minor axis) The upper limit of (average particle diameter) is 3 μm. Here, the crystal grain is a base crystal grain surrounded by a grain boundary having a crystal orientation difference of 15 ° or more. On the other hand, when the average particle size of the long axis of the dispersed particles is larger than 0.3 μm, the particle dispersion strengthening can hardly be expected, and the steel having 1.2 GPa or more is highly likely to significantly deteriorate the toughness. Therefore, it is desirable that the average particle diameter of the major axis is 0.3 μm or less.

  In particular, the effect of refining crystal grains becomes particularly remarkable in the region where the average crystal grain size is 1 μm or less, and the particle dispersion strengthening by the Orowan mechanism is in the region where the average particle size is 0.1 μm or less. Therefore, in order to effectively utilize the strengthening by crystal fiber formation and the particle dispersion strengthening, it is effective to further reduce the minor axis average particle diameter of the fibrous crystal to 1 μm or less, further 0.5 μm or less. . The average particle diameter of the major axis of the second phase dispersed particles is also preferably 0.1 μm or less, more preferably 0.05 μm or less, depending on the refinement of the matrix structure.

  In such warm-worked steel materials, strengthening mechanisms such as solid solution strengthening and dislocation strengthening can be added in addition to the strengthening mechanism described above, and a simple addition of the strengthening mechanism due to the effect of superimposing these strengthening mechanisms. In this way, highly functional materials that cannot be predicted have been obtained.

  Such a fine fiber structure can be formed by warm forming of a plate material, a rod wire, a bolt thread portion, and the like. In particular, even when the amount of accumulated strain is small, a fiber structure can be formed in a surface layer portion or the like that has undergone strong local deformation, and the characteristics of various parts and desired portions can be greatly improved.

  Embodiments of the present invention will be described in more detail below with reference to the accompanying drawings. Of course, the present invention is not limited to the following examples, and it goes without saying that various aspects are possible in detail.

  Table 1 shows steel components (A to K, M, N and O) within the range of the present invention and steel components (L) outside the range. In the examples, carbide was used as the second phase dispersed particles. Table 2 shows the volume ratios of metal carbide and cementite that can be dispersed as second phase dispersed particles in the steel having the composition shown in Table 1. The steels of the examples cover martensitic steels ranging from SCM435 to 2GPa grade secondary hardened steel, excluding Co-added maraging steel.

  There are various stoichiometric carbides in the actual steel depending on the steel components and heat treatment conditions. For this reason, it is difficult to strictly measure the volume fraction of the second phase dispersed particles by chemical analysis or structural observation, which is not practical. Therefore, the inventors calculated the volume fraction of carbides by calculation from the well-known theoretical density of carbides obtained by structural analysis of carbides (Tokyo Chemical Doujin, Chemical Dictionary, (1989), P.1361-1363). Asked. Table 3 shows approximate equations for calculation.

  In the calculation, it was assumed that carbides are formed by combining with carbon in the order of alloy elements having a strong carbide forming ability (Nb> Mo> Cr> Fe, etc.). It is well known that Nb and Mo are elements that make it easy to produce unique carbides in steel and are difficult to dissolve in cementite, and precipitation of NbC and Mo2C was assumed. However, about G steel and L steel, since 0.002 mass% Mo is the quantity which can fully dissolve in cementite, it excluded from the estimation of the volume fraction of Mo carbide. As for Cr, when a large amount of Cr is added, carbides such as M23C6 and M7C3 having a high Cr concentration are formed. However, with the addition amount of this example, Cr forms a solid solution in cementite to form these alloy carbides. The possibility of doing is low. Therefore, the volume fraction of Cr alloy carbide was excluded from the estimation.

The most important thing here is that in medium-carbon low alloy steel, the amount of dispersion of carbides, which are dispersion strengthened particles, depends on the amount of carbon, especially when there is no possibility of forming metal carbide with a sufficiently high density relative to cementite. In addition, when the amount of the element forming the metal carbide is small, the amount of the second phase dispersed particles is almost determined by the amount of cementite. That is, as shown in Table 2, the total amount of the volume fraction of the second phase sufficiently exceeds 7 × 10 ⁻³ in the steel having the C content of 0.2 mass% or more used in the examples of Table 1.

  1, 2, and 3 exemplify the heat treatment process applied in the embodiment. This process basically consists of (1) solution heat treatment and processing for reducing coarse undissolved carbide, and (2) tempered martensite or bainite structure as the structure of the steel for warm working of the present invention. It consists of quenching and tempering to obtain, and (3) warm working that also serves to shape the parts. In the thermomechanical processing pattern 1 of FIG. 1, the reverse transformation austenite grains are refined by austenitizing at a low temperature following the solution heat treatment, and in the pattern 2 of FIG. 2, recrystallization obtained by hot working subsequent to the solution heat treatment. Considering quenching from austenite and unrecrystallized austenite (elongated austenite) structure obtained by warm working. FIG. 3 shows a quenching process from a processed austenite (elongated austenite) structure by ausforming treatment in a metastable austenite region. In these thermomechanical processes, the finer the crystal grains, the finer the microstructure can be obtained by warm processing with a smaller cumulative strain amount. In particular, the microstructure before processing for efficiently developing the fiber structure is not fine. It is most effective to use martensite obtained from recrystallized austenite (elongated austenite) as a pre-structure.

  First, about 40 mm square × 120 mm long square cut out from a hot-rolled steel sheet or forged material is subjected to quenching treatment in the heat treatment patterns 1, 2 and 3 to obtain a martensite single structure close to almost 100% by volume. It was. This corresponds to an example of the steel for warm working according to the present invention. The square was then heated to a predetermined temperature in 0.5 hours and tempered, and then subjected to warm rolling using a grooved roll to a predetermined area reduction rate to impart strain and air-cooled.

  The structure of the obtained steel material was observed by polishing a cross section parallel to the rolling (RD) direction using an optical microscope, a transmission electron microscope (TEM), and an FE-SEM and EBSP analyzer. The prior austenite grain size was determined in accordance with a comparison method or cutting method defined in JIS G 0552 by corroding the polished surface with an aqueous picric acid alcohol solution to reveal prior austenite grain boundaries. The average particle size of the second phase dispersed particles is measured with three or more fields of view at a magnification of 10,000 to 100,000 times using TEM or SEM, and the major axis length of 250 or more particles in total is measured. And asked. When some particles were coalesced and aggregated, it was regarded as one particle. The maximum particle size was made to correspond to the length of the long axis of the largest carbide among the measured carbides. The average grain size of the short axis and long axis of the elongated grains in the fiber structure was measured by the cutting method by the EBSP analysis by measuring the average section length of the short axis and long axis of the elongated grains having a crystal orientation difference of 15 ° or more. (See FIG. 5).

  The hardness of the obtained steel material was measured at a load of 20 kg and a holding time of 15 s using a Vickers hardness tester in accordance with a test method defined in JIS Z 2244.

  The tensile test conforms to the test method specified in JIS Z 2241. 1) Parallel part diameter 3.5 mm, length 24.5 mm, distance between scores 17.5 mm, or 6 mm, length 42 mm, distance between scores A 30 mm JIS No. 14 A proportional test piece, or 2) a JIS No. 4 subsize test piece having a parallel part diameter of 10 mm, a length of 45 mm, and a distance between marks of 35 mm, was performed at room temperature using an Instron type tensile tester. The crosshead speeds of 1) JIS No. 14A and 2) JIS No. 4 were 0.5 mm / min and 10 mm / min, respectively, and the elongation was measured up to the breakage by attaching an extensometer to the test piece.

  The impact test is a U-notch or V-notch test piece having a length of 55 mm and a height and width of 10 mm produced by cutting a steel material having a cross-sectional area of 1.8 cm 2 or more in accordance with a test method defined in JIS Z 2242. I followed.

  Hydrogen embrittlement characteristics were evaluated at room temperature at a crosshead speed of 0.005 mm / min using a low strain rate tensile tester for a notched specimen with a diameter of 10 mm, a notched bottom diameter of 6 mm, and a stress concentration factor of 4.9. did. In the hydrogen embrittlement test, the average amount of hydrogen in the test piece was changed by cathode charging for 72 hours with the charge liquid and current density changed, and Cd plating was applied to prevent the hydrogen in the test piece from being dissipated. The above test was conducted. Hydrogen analysis was performed on the sample from which Cd plating had been removed by a temperature programmed desorption hydrogen analysis method using a quadrupole mass spectrometer, and the hydrogen released up to 300 ° C. was defined as diffusible hydrogen. .

  Table 4 summarizes the results of evaluating the manufacturing conditions and structure of the steel for warm working, the quenching and tempering conditions of the unprocessed material and their hardness, and the suitability as the warm working steel of the present invention.

  FIG. 4 shows the relationship between T (logt + 20) = λ and the hardness of the tempered martensitic steel as it is.

In the comparative steel L, the volume fraction of cementite is 33 × 10 ⁻³ , but the cementite is not thermally stable because it does not contain the alloying elements specified in the present invention, and is easily grown by heating. Resulting in. Therefore, in the tempering treatment of λ = 1.4 × 10 ⁴ or more, the hardness of the L steel is less than H = (5.2-1.2 × 10 ⁻⁴ λ) indicated by a broken line in the figure, and 350 HV3.7 × 10 ² cannot be achieved with L steel by warm working above ℃.

FIG. 5 shows the result of steel I being gamma-ized at 11.0 × 10 ² ° C., water-cooling, tempering at 5.0 × 10 ² ° C. for 1.5 hours, and then hot groove rolling. An example of analyzing the structure of the obtained material is shown. The accumulated strain applied at this time is 2.4, and the hardness is HV 3.7 × 10 ² . As can be seen from the EBSP analysis diagram (a) and the TEM photograph (b) of the Bcc phase, an ultrafine fiber structure in which spherical carbides are dispersed in the ferrite phase base elongated in a fibrous form is obtained. As a result of measuring the minor axis average grain size of the crystal grains having a crystal orientation difference of 15 ° or more by the EBSP analysis (c), the minor axis average grain diameter of the elongated crystal grains is 0.3 μm. there were. However, in this steel, the fiber structure was complicatedly developed, and the average particle diameter of the major axis could not be measured. On the other hand, as a result of measuring the particle diameter (major axis length) of 287 carbides by TEM, the average particle diameter of the carbides was 0.06 μm and the maximum diameter was 0.2 μm (d).

  From the inverse pole figure regarding the rolling direction (RD), it can be seen that the <011> // RD texture is a developed fiber structure. Similar textures were formed for other developed steels. Since the cleavage plane of Bcc iron is {100}, the formation of such <011> fiber structure is extremely effective for fracture due to tensile deformation in the fiber axis direction or bending deformation that receives a bending moment along the fiber direction. I think there is.

  Table 5 shows the relationship between warm working conditions and the structure and hardness of the obtained warm worked material. In the table, T and t are the processing temperature and processing time shown in FIGS. 1 to 3, respectively.

The hardness of the work material greatly depends on the temper softening resistance. When compared with the same λ = T (logt + 20), a steel having a higher temper softening resistance can obtain a higher work material. In particular, in a processed material of HV 4.0 × 10 ² or more, the base structure is ultra-fine to an average width of 0.5 μm or less. In the processed material of HV 4.0 × 10 ² or more, since extremely fine particles are densely dispersed, the average carbide particle diameter could not be determined strictly, but relatively particles such as I steel in FIG. When compared with a larger one, it was possible to determine that it was less than 0.1 μm.

However, even in the case of steel for warm working with high resistance to temper softening, carbide particles and the like easily grow during processing in a high temperature region of 700 ° C., so that HV 3.7 × 10 It becomes difficult to obtain ^two or more warm processed materials (Comparative Examples 3, 4, 5, and 7). Therefore, it is desirable that processing in such a high temperature range be performed by combining heating and processing for a short time so that carbides and the like do not grow using high-frequency heating, for example. Further, in the processing in a high temperature region around 700 ° C., grain growth is likely to occur, so the proportion of crystal grains having a small aspect and a relatively large grain size increases. As a result, the degree of extension becomes small. For example, in Comparative Examples 4, 5, and 7, the degree of extension was measured as 6, 2, and 4, respectively. Although the degree of extension could not be measured for the examples, the degree of extension could be determined to be 6 or more in comparison between Comparative Example 7 and the tissue.

  Tables 6 and 7 summarize examples and comparative examples of mechanical properties. In the table, UE and VE are absorbed energy of U-notch and V-notch test pieces, respectively.

  The steel having the composition shown here is a steel for warm working that has been subjected to appropriate alloy design and heat treatment so that the second phase dispersed particles are finely dispersed. 16 or more tensile strength x total elongation balance. However, when compared with the same composition, a larger tensile strength × total elongation balance is obtained with the developed steel subjected to warm working than the comparative example. Moreover, even when carbon of about 0.2 mass% is added, if an alloying element such as Mo is appropriately blended, it is noticed that a 1.5 GPa grade almost equal to the quenching hardness can be obtained, and that the ductility is excellent. (A, D, E steel). Furthermore, as for the O steel, a 2GPa grade ultra high strength steel was obtained.

  On the other hand, it can be seen from the shock absorption energy that the developed steel has much better toughness than conventional high-strength steels up to low temperatures.

FIG. 6 summarizes the relationship between the tensile strength and the impact value at room temperature (U-notch test piece). In addition, the data of the steel for machine structural standards (New Japan Casting Forging Association: collection of steel materials for machine structural materials for field use (1995)) standardized by JIS are also shown in the figure. In the conventional steel, the impact value is significantly reduced in the strength range of 1.2 GPa or more, and in the strength of 1.5 GPa or more, it is 70 J / cm ² or less, whereas in the steel of the present invention, even in the strength of 1.5 GPa or more. An extremely high impact value of 150 J / cm ² or more is exhibited.

  FIG. 7 summarizes the relationship between tensile strength and absorbed energy at room temperature (V-notch test piece). In the figure, data for steel for machine structural use as defined by JIS (Kinki Giken Fatigue Data Sheet Material 5) is also shown. The present invention is excellent in toughness in a high strength region as compared with conventional ausformed steel, fine grain steel, maraging steel and the like.

  FIG. 8 shows the relationship between test temperature and absorbed energy. For example, from Example 1, Comparative Example 8, Examples 3, 5, and Comparative Example 10, it can be confirmed that a material having high absorbed energy can be obtained by processing. In particular, it is noteworthy that some developed steels exhibit not only higher absorption energy near room temperature, but also a unique temperature dependence that shows a maximum value in the low temperature range. For example, in steel A of Example 1 and steel B of Example 3, a peak is observed at around −40 ° C., and in steel F of Example 11, a peak of around −100 ° C. is observed, and there are some unbroken parts in the peak temperature range. did. As shown in FIG. 9, the developed steel is characteristic in that it has a fiber shape as if the fracture surface is a bamboo fold. A similar phenomenon is observed when an ausformed 0.2 mass% C-3 mass% Ni-3% Mo steel (tensile strength: 1.6 GPa) is tested at around 200 ° C. The absorbed energy is reduced to about 33J (Non-patent Document 16). In addition, even when a steel obtained by modified ausfoam treatment of 0.5 mass% C-0.9 mass% Mn-0.8 mass% Cr steel (5150 steel) is subjected to an impact test, fibrous fracture occurs and the toughness is improved. However, the maximum absorbed energy at room temperature is about 90 J at an intensity level of 1.5 GPa (Non-patent Document 17). Therefore, in the tensile strength of 1.2 GPa or more, the absorbed energy near normal temperature is much higher than that of the existing ausfoam steel as in the newly developed steel, and the absorbed energy in the low temperature range of -40 ° C or lower. It is a notable finding that has the highest value.

  As described above, the mechanical properties, particularly high impact properties, of the excellent developed steel are largely attributed to the ultrafine <011> fiber structure that has been densely developed by warm working of the particle-dispersed multiphase structure.

  FIG. 10 shows the relationship between the hardness of the warm processed material and the aging temperature. In a warm processed material to which a secondary hardening element such as Mo is added, it is possible to maintain the hardness up to a high temperature by aging treatment or to increase the strength as compared with the warm processing.

  FIG. 11 shows an example in which the ultrafine fiber structure formed in the center of the plate material by warm rolling of N steel at 650 ° C. is observed by SEM.

  FIG. 12 shows an example in which an ultrafine fiber structure formed in the surface layer portion of a bar material that is locally strongly deformed is observed with an SEM.

  Table 8 shows the results of the hydrogen embrittlement resistance test.

  Here, a notched tensile test is performed on a steel charged with a hydrogen amount of about 0.3 mass ppm, and the notched tensile strength at that time is 0.7 times the tensile strength of a smooth tensile test piece not charged with hydrogen. The hydrogen embrittlement resistance was evaluated based on the result.

  The developed steel satisfies this condition even at a high strength level with a tensile strength of 1.6 GPa or more, and can be judged to be excellent in hydrogen embrittlement resistance. In addition, the comparative example 14 is the steel for high-strength mechanical structure excellent in the delayed fracture invented in Japanese Patent Application No. 2001-264399.

  Even for high-strength steels that have been made into a multi-phase by fine dispersion of a small amount of second-phase dispersed particles, especially ultra-high-strength steels that are difficult to soften and difficult to form, deformation resistance is reduced and cracks occur in the material By applying a predetermined deformation in a temperature range that does not cause molding into a predetermined shape (thin plate, thick plate, bar wire, part), the conventional spheroidizing annealing and quenching and tempering processes after forming the part are omitted. Provided is a high-strength steel and a member in which an ultrafine multiphase structure is developed into a fibrous shape, and ductility having a trade-off balance with high strength, particularly toughness and hydrogen embrittlement resistance are greatly improved.

  These are useful as steels or members that are processed into various structures and automobile parts.

FIG. 1 is a diagram showing an example of a heat treatment pattern. FIG. 2 is a diagram showing an example of a heat treatment pattern. FIG. 3 is a view showing an example of the heat treatment pattern. FIG. 4 is a diagram illustrating the relationship between tempering hardness and T (logt + 20) = λ, where T is a tempering temperature (K) and t is a tempering time (hr). FIG. 5 is a diagram illustrating a 500 ° C. warm-worked structure (ultrafine fiber structure). FIG. 6 is a diagram illustrating the relationship between tensile strength and impact value (U notch). FIG. 7 is a diagram illustrating the relationship between tensile strength and absorbed energy (V notch). FIG. 8 is a diagram illustrating the relationship between absorbed energy and test temperature. FIG. 9 is a photographic view showing an example of a fracture form of B steel subjected to Charpy impact test (U notch). FIG. 10 is a diagram illustrating the relationship between the hardness of the warm processed material and the aging temperature. FIG. 11 is a diagram illustrating an ultrafine fiber structure formed in the center of the plate material. FIG. 12 is a diagram illustrating an ultrafine fiber structure formed in the surface layer portion of the bar.

Claims (13)

The chemical composition is
C: 0.70 mass% or less,
Si: 0.05 mass% or more and 2.5 mass% or less,
Mn: 0.05 mass% or more and 3.0 mass% or less,
Cr: 0.01 mass% or more and 2.01 mass% or less,
Al: 0.5 mass% or less,
O: 0.3 mass% or less,
N: 0.3 mass% or less,
Mo: 0.002 mass% or more and 5.0 mass% or less,
The balance is steel that is Fe and inevitable impurities,
The steel material is composed of a particle-dispersed fibrous crystal grain structure exhibiting a <011> // RD (rolling direction) texture, and the average grain diameter of the short axis of the fibrous ferrite crystal forming the base structure is 3 μm or less. The phase particles are finely dispersed in the matrix structure at a volume ratio of 7 × 10 ⁻³ or more and 12 × 10 ⁻² or less, and the Vickers hardness of the steel material at room temperature is HV3.7 × 10 ² or more. Steel material.   2. The steel material according to claim 1, wherein the base structure is made of a fibrous ferrite crystal having a minor axis average particle diameter of 1 μm or less.   3. The steel material according to claim 1, wherein the base structure is made of a fibrous ferrite crystal having a minor axis average particle diameter of 0.5 μm or less.   The steel material according to any one of claims 1 to 3, wherein the average particle diameter of the major axis of the second phase dispersed particles is 0.1 µm or less. The steel material according to any one of claims 1 to 4 ,
W: 5.0 mass% or less,
V: 5.0 mass% or less,
Ti: 3.0 mass% or less,
Nb: 1.0 mass% or less,
Ta: Steel material characterized by further containing one or more selected from the group consisting of 1.0 mass% or less. The steel material according to any one of claims 1 to 5 ,
Ni: 0.05 mass% or more and 9 mass% or less ,
Cu: Steel material characterized by further containing 1 type or 2 types of 2.0 mass% or less. It is steel for warm processing for creating the steel materials of any one of Claim 1 to 6 by warm processing, Comprising: In the predetermined temperature range of 350 degreeC or more and Ac1 point or less, following formula (1) The second phase particles are generated by performing heat treatment of any one of annealing , tempering, or aging treatment under the condition that the parameter λ represented is 1.4 × 10 ⁴ or more, and at room temperature after this heat treatment, The volume ratio of the second phase particles is 7 × 10 ⁻³ or more and 12 × 10 ⁻² or less, and the Vickers hardness (HV) of the steel material is a hardness H or more of the following formula (2). Inter-working steel.
λ = T (logt + 20) (T; temperature (K), t; time (hr)) (1)
H = (5.2-1.2 × 10 ⁻⁴ λ) × 10 ² (2) The warm-working steel according to claim 7 , wherein 80% by volume or more of the base structure is a single structure of martensite or bainite, or a mixed structure thereof. A steel plate material obtained by warm-working the warm-working steel according to claim 7 or 8 , wherein at least the surface layer portion has a particle-dispersed type according to any one of claims 1 to 6. A steel plate material in which a fibrous crystal structure is generated. A steel rod obtained by warm-working the warm-working steel according to claim 7 or 8 , wherein the particle dispersion according to any one of claims 1 to 6 is at least on a surface layer portion thereof. A steel bar characterized in that a type fibrous crystal structure is generated. A steel bolt obtained by warm-working the warm-working steel according to claim 7 or 8 , wherein the particles according to any one of claims 1 to 6 are at least on a surface layer portion of a screw portion. A steel bolt, wherein a dispersed fibrous crystal structure is generated. A method for producing the steel material according to any one of claims 1 to 6 , the steel plate material according to claim 9 , the steel bar material according to claim 10 , or the steel bolt according to claim 11. In the warm working steel according to claim 7 or 8 , the parameter λ represented by the following formula (1) is 1.4 × 10 ⁴ or more in a predetermined temperature range of 350 ° C. or more and Ac1 point or less. A warm working method characterized in that a predetermined shape is obtained by warm working with a strain of 0.7 or more under the following conditions.

λ = T (logt + 20) (T; temperature (K), t; time (hr)) (1) A steel product obtained by cutting a steel material, wherein the steel material is a steel material according to any one of claims 1 to 6 , a steel plate material according to claim 9 , and a steel rod according to claim 10. A steel processed product characterized by being either a material or a steel bolt according to claim 11 .