JP7553065B2 - Ultrafine ferrite-cementite structure steel, ultrafine ferrite-austenite structure steel, ultrafine martensite structure steel, and method for producing ultrafine martensite-austenite structure steel - Google Patents

Description

Article 30, paragraph 2 of the Patent Act applies. Distribution date: September 1, 2018 Publication: 176th Autumn Meeting of the Iron and Steel Institute of Japan (Tohoku University) Lecture Abstracts CD-ROM version Issue date: December 14, 2018 Publication: Tetsu to Hagane, Vol. 105, No. 2, pp. 197-206 Distribution date: March 1, 2019 Publication: 177th Spring Meeting of the Iron and Steel Institute of Japan (Tokyo Denki University) Lecture Abstracts CD-ROM version Distribution date: August 26, 2019 Publication: 178th Autumn Meeting of the Iron and Steel Institute of Japan (Okayama University) Lecture Abstracts CD-ROM version

The present invention relates to a method for producing ultrafine ferrite-cementite structured steel, ultrafine ferrite-austenite structured steel, ultrafine martensite structured steel, and ultrafine martensite-austenite structured steel, which have high strength and high ductility and are suitable for use in structures such as buildings and bridges, automobile chassis steel, and machine gears and other parts.

In recent years, from the perspective of protecting the global environment, improving fuel efficiency by reducing the weight of automobile bodies and collision safety from the perspective of protecting passengers have become more important than ever. It is possible to strengthen a car body by increasing the thickness of the steel plate or adding reinforcing members, but this increases the weight of the car body. One technology that meets these conflicting demands is the development of high-strength steel plate, or so-called hi-tensile steel. However, increasing the strength of steel plate comes with the issue of deterioration of formability, so it is necessary to develop high-strength, highly formable steel plate with high formability according to the parts to which it is applied. Therefore, hi-tensile steel, which uses a retained austenite-type composite structure, has attracted attention as a way to achieve both high strength and high ductility.

When this high-tensile steel containing metastable austenite is deformed, the austenite hardens during deformation due to the work-induced transformation of the austenite to martensite, suppressing localization of deformation and allowing the deformation to transfer to untransformed austenite, resulting in high elongation. Since plastic deformation is promoted by work-induced martensite transformation, this phenomenon is called the TRIP (Transformation Induced Plasticity) phenomenon, and this retained austenite steel is called TRIP steel. Currently, the target for high-tensile steel for next-generation automobiles is said to be tensile strength TS x total elongation T.EI ≧ 30,000 MPa%, and material development to achieve this goal is underway.

In light of this background, it has been reported that low-C, medium-Mn steel with a ferritic-austenite structure and a composition of 0.1%C-2%Si-5%Mn can be manufactured by extremely simple interphase annealing compared to conventional TRIP steel, and also exhibits high strength (1200 MPa) and high ductility (T.EI = 25%) (see, for example, non-patent documents 1 and 2).

S. Torizuka and T. Hanamura: Bull. Iron Steel Inst. Jpn., 17(2012), 852. T. Hanamura, S. Torizuka, A. Sunahara, M. Imagumbai and H. Takechi: ISIJ Int., 51(2011), 685.

However, in ferrite-austenite steel with the above composition, the annealing time required to obtain the optimal austenite fraction (20-30%) is long, at about 1 hour, making it difficult to manufacture by continuous annealing. This makes practical use difficult. Therefore, it is necessary to develop a method to efficiently obtain the desired ferrite-austenite steel in a short time of about 10 minutes.

As a result of intensive research, the inventors discovered that by using ultrafine ferrite-cementite structure steel as the precursor structure of ferrite-austenite structure steel, ultrafine ferrite-austenite structure steel can be obtained in an annealing time of just 10 minutes, one-sixth the time required by conventional methods. Furthermore, they discovered that by adjusting the content of Mn among the metal elements that are the main raw materials, it is possible to produce ultrafine equiaxed martensite with high strength and high ductility by quenching at a relatively low temperature of 750°C.

The present inventors have studied this mechanism in detail and found that in the steel with an ultrafine ferrite-cementite structure as a precursor structure, cementite that is originally in the form of Fe ₃ C has a part (50 wt%) of Fe replaced by Mn, such as (Fe ₅ ,Mn ₅ ) ₃ C. In this way, it has been found that the presence of a high concentration of Mn in cementite increases the rate of reverse transformation from ferrite to austenite and allows the annealing temperature to be lowered.

The present invention was made based on the above findings, and aims to provide high-strength, high-ductility ultrafine ferrite-cementite structure steel, ferrite-austenite structure steel, ultrafine martensite structure steel, and ultrafine martensite-austenite structure steel that have sufficient mechanical strength and ductility even when used in parts such as buildings and bridges, automobile suspension steel, and machine gears, and to provide a manufacturing method that can produce these at relatively low temperatures in an extremely short time.

According to the present invention, in order to solve the above problems, the following technical means or techniques are provided.
[1] A steel with an ultrafine ferrite-cementite structure, the contents of C, Si and Mn being 0.05 wt%≦C≦0.3 wt%, 0.5≦Si≦2 wt%, 3 wt%≦Mn≦10 wt%, respectively, with the balance being Fe and unavoidable impurities, and the ferrite grain size being 2.0 μm or less, wherein Mn is dissolved in the cementite, and the value of x in the chemical formula of the cementite (Fe1-x, Mnx)3C is 0.3 or more and less than 1.
[2] An ultrafine-grained ferrite-austenite structure steel having a structure before annealing in a dual-phase region, the ultrafine-grained ferrite-cementite structure steel of the first invention being the above-mentioned, characterized in that the ultrafine-grained ferrite-austenite structure steel has a tensile strength of 1400 MPa or more and a total elongation of 25% or more.
[3] The second aspect of the present invention provides a high-strength, high-ductility, ultrafine-grained ferrite-austenite steel, characterized in that the steel has a tensile strength of 1500 MPa or more and a total elongation of 30% or more.
[4] The ultrafine-grained ferrite-austenite structure steel according to the third invention, characterized in that the austenite volume fraction is 40% or more and V _γ xCin _γ ≧0.07.
[5] An ultrafine martensite structure steel having a structure before austenitic region annealing of the ultrafine ferrite-cementite structure steel of the first invention, characterized in that the ultrafine martensite structure steel has a tensile strength of 1500 MPa or more and a total elongation of 15% or more.
[6] A martensite-austenite structure steel having a structure before annealing in the austenite region and two-phase region, the ultrafine-grained ferrite-cementite structure steel of the first invention being the above-mentioned, characterized in that the ultrafine-grained martensite-austenite structure steel has a tensile strength of 1500 MPa or more and a total elongation of 20% or more.
[7] A method for producing a steel with an ultrafine ferrite-cementite structure according to the first invention, comprising the steps of: subjecting a raw material having the above-mentioned chemical components to warm rolling at a rolling reduction rate of 50% or more in a temperature range of 300 to 600°C to form a fine ferrite-cementite structure; and concentrating Mn in the cementite.
[8] A method for producing an ultrafine-grained ferrite-cementite structure steel produced by the seventh invention, characterized in that the steel is held at a temperature of 625 to 750°C for 180 to 1200 seconds, and then water-cooled or air-cooled.
[9] A method for producing ultrafine martensite structure steel or martensite-austenite structure steel, characterized in that the ultrafine ferrite-cementite structure steel of the seventh invention is held at a temperature range of 625 to 800°C for 1 to 600 seconds, and then water-cooled or air-cooled to produce a structure which is equiaxed martensite having a single block structure or an ultrafine equiaxed martensite-austenite structure steel having a single block structure.

The present invention employs the above-mentioned configuration, and therefore has the excellent effect of providing sufficient mechanical strength and ductility even when used in parts such as buildings and bridges, automobile suspension steel, and machine gears. It also has the excellent effect of being able to produce ultrafine ferrite-cementite structure steel at a relatively low temperature in an extremely short time.

(a) is an SEM image of a fine ferrite-cementite (α+θ) structure, which is a structure before annealing in the two-phase region, and (b) is an SEM image of a martensite (μ) structure. FIG. 2 is a diagram showing the relationship between the two-phase annealing time and the structure. FIG. 2 is a graph showing the change in austenite volume fraction with the change in two-phase annealing temperature and time. 1(a), (b), and (c) are schematic diagrams of the process of austenite (γ) formation by reverse transformation. (a) shows a fine ferrite-cementite (α+θ) structure before annealing, and (b) shows an SEM image of the structure obtained when the ferrite-cementite (α+θ) structure before annealing and the martensite (M) structure were heated to 675°C and immediately water-cooled without being held at that temperature. (a) shows a fine ferrite-cementite (α+θ) structure before annealing, and (b) shows a TEM image of the structure obtained when the ferrite-cementite (α+θ) structure before annealing and the martensite (M) structure were heated to 675°C and immediately water-cooled without being held at that temperature. FIG. 1 shows the results of TEM-EDS analysis. FIG. 1 is a diagram showing the results of Mn concentration mapping by a field emission electron probe microanalyzer (FE-EPMA) for a ferrite-austenite (α+γ) structure after dual-phase annealing. FIG. 3 is a diagram showing nominal stress-nominal strain curves obtained in a round bar tensile test of the four ferrite-austenite (α+γ) structures shown in FIG. 2. This figure shows the nominal stress-nominal strain curves of plate-shaped tensile test pieces of ferrite-austenite (α+γ) structure steel, which had a ferrite-cementite (α+θ) structure and a martensite (M) structure before annealing and were annealed at 675°C for 3600 s, and also shows the change in austenite (γ) volume fraction during tensile deformation. This is a diagram plotting the relationship between the strength and ductility balance of ferrite-austenite (α+γ) structure steel and conventional TRIP steel, based on the results of the plate tensile test shown in FIG. This is a graph showing nominal stress-nominal strain curves of ferritic-austenitic steel with a 2% Si-5% Mn composition, annealed at 675° C. for 1 hour with varying C contents of 0.075, 0.1, 0.15, 0.2, and 0.3 wt. This is a graph showing nominal stress-nominal strain curves of ferritic-austenitic steel with a 2% Si-5% Mn composition, annealed at 700° C. for 1 hour with varying C contents of 0.075, 0.1, 0.15, 0.2, and 0.3 wt. FIG. 1 is a graph showing the relationship between uniform elongation and V _γ xCin _γ of an ultrafine-grained ferrite-austenite steel having a low carbon-2% Si-5% Mn composition. FIG. 1 is a graph showing the relationship between tensile strength and austenite volume fraction _Vγ of ultrafine-grained ferrite-austenite steel with a low carbon-2% Si-5% Mn composition. FIG. 1 is a diagram showing a nominal stress-nominal strain curve of a ferrite-austenite structure steel having an ultrafine ferrite-cementite structure steel with a composition of 0.15C-2%Si-5%Mn as a structure before annealing.

The present invention will be described in detail below with reference to the embodiments. In this specification, the steel structures described as ultrafine ferrite-cementite steel, ultrafine ferrite-austenite steel, and ultrafine martensite-austenite steel refer to dual-phase steel with the structure described before and after the hyphen.

The high-strength, high-ductility ultrafine ferrite-cementite structure steel of the present invention is a high-strength, high-ductility ultrafine ferrite-cementite structure steel with the C, Si and Mn contents being 0.05 wt%≦C≦0.3 wt%, 0.5≦Si≦2 wt%, 3 wt%≦Mn≦10 wt%, respectively, with the remainder being Fe and unavoidable impurities, and with a ferrite grain size of 2.0 μm or less, characterized in that Mn is dissolved in the cementite, and the value of x in the chemical formula of the cementite (Fe1-x, Mnx)3C is 0.3 or more and less than 1.

In the ferrite-cementite structure steel of the present invention, the C content is 0.05 wt%≦C≦0.3 wt%. C is necessary to ensure hardenability and tensile strength, but if it is less than 0.05 wt%, the tensile strength of the steel may not be fully satisfied, and if it exceeds 0.3 wt%, the ductility and toughness of the steel may decrease, and the weldability may decrease.

In the ferrite-cementite structure steel of the present invention, the Si content is 0.5 wt%≦Si≦2 wt%. Si is a substitutional element that greatly hardens the material, and is an effective element for improving the strength of steel. However, if the Si content is excessively high, a lot of Si scale will be generated during heating during hot working, which will result in extra costs for scale removal and will make surface defects more likely to occur due to the scale. Therefore, it is desirable to set the upper limit of Si to 2 wt%.

In the ferrite-cementite structure steel of the present invention, the Mn content is 3 wt%≦Mn≦10 wt%. Mn is an austenite-forming element that improves hardenability and forms martensite from fine γ grains. Furthermore, if the Mn content is within the above range, even with annealing at a relatively low temperature and for a short time, Mn is dispersed and dissolved in the cementite formed during warm rolling at a high concentration where the value of x (weight ratio) in (Fe1-x, Mnx)3C is 0.3 or more and less than 1, for example, a concentration of about 0.3 to 0.55, and it is believed that the high concentration of Mn effectively acts as a preferential nucleation site for reverse transformed austenite. As a result, highly stable austenite can be formed in a short time.

On the other hand, if the Mn concentration becomes too high, the segregation of Mn in the steel during solidification becomes excessive, impairing the uniformity inside the material. In addition, surface cracks become more likely to occur during the hot working process in the material preparation process. Therefore, it is desirable to set the upper limit of Mn to 10 wt%.

The microstructure of the ultrafine ferrite-cementite steel of the present invention has a ferrite grain size of 2.0 μm or less, more preferably 1.0 μm. If the ferrite grain size exceeds 2.0 μm, the yield point decreases, which is undesirable.

The ultrafine ferrite-austenite structure steel of the present invention is an ultrafine ferrite-austenite structure steel in which the ultrafine ferrite-cementite structure steel is used as a pre-annealing structure in the dual-phase region, and has a tensile strength of 1400 MPa or more and a total elongation of 25% or more. When the ultrafine ferrite-cementite structure steel is used as a pre-annealing structure in the dual-phase region and then annealed, a higher austenite volume fraction is obtained than in conventional TRIP steel, and although the amount of solute C in the austenite is low, the effect of the high concentration of solute Mn stabilizes the austenite and high ductility can be obtained.

When the ultrafine ferrite-cementite structure steel of the present invention is subjected to a two-phase annealing process before annealing, a high-strength, high-ductility ultrafine ferrite-austenite structure steel can be obtained, characterized by a tensile strength of 1500 MPa or more and a total elongation of 30% or more. For example, by annealing for an extremely short period of time, such as about 10 minutes, it is possible to obtain an ultrafine ferrite-austenite structure steel that has higher yield stress, tensile strength and total elongation than conventional TRIP steel and has an extremely good balance of strength and ductility.

Furthermore, in the ultrafine-grained ferrite-austenite steel of the present invention, it is preferable that V _γ xCin _γ ≧0.07 when the austenite volume fraction is 40% or more, and more preferably, V _γ xCin _γ ≧0.14.

The ultrafine martensite structure steel of the present invention is a martensite structure steel in which the ultrafine ferrite-cementite structure steel is used as the structure before annealing in the austenite region, and is preferably considered to have a tensile strength of 1500 MPa or more and a total elongation of 15% or more.

The ultrafine martensite-austenite structure steel is preferably considered to have a tensile strength of 1500 MPa or more and a total elongation of 20% or more.

Furthermore, the martensite-austenite structure steel of the present invention is a martensite-austenite structure steel in which the high-strength, high-ductility ultrafine ferrite-cementite structure steel is used as the structure before annealing in the dual-phase region, and is preferably considered to have a tensile strength of 1400 MPa or more and a total elongation of 25% or more. Desirably, it is preferably considered to have a tensile strength of 1500 MPa or more and a total elongation of 30% or more.

Martensite has a complex hierarchical structure made up of four components. Prior austenite (γ) grains, which are 30 μm in size, are packed with packets that are several μm in size, and these pockets are made up of long, thin blocks that are about 1 μm wide. These blocks are further made up of laths. In other words, the four components of prior austenite grains, packets, blocks, and laths are stacked on top of each other. This creates a complex hierarchical structure in which carbide particles several to tens of nm in size are dispersed at the grain boundaries and boundaries between these four components, as well as within the grains.

Normally, one prior austenite (γ) grain in martensite has multiple packets, and within each packet, blocks with various orientations (variants) are transformed into and formed. The block boundaries are high-angle grain boundaries. However, as the γ grain size is refined, single-packet multi-blocks may form, and when refined even further, single-block (single-variant) martensite may form. In other words, as prior austenite is refined, the number of blocks that can form within it decreases. Eventually, only one block can form. It is as if one prior austenite grain has seemingly transformed into one ferrite grain. However, in general, refinement of the γ grain size reduces hardenability and does not result in martensite.

However, in the present invention, by adding an appropriate amount of Mn and forming the structure before dual-phase annealing into ultrafine ferrite-cementite, it is possible to make the prior gamma grain size 2.0 μm or less, and by further adding an appropriate amount of Mn, it is possible to make the structure into equiaxed martensite with a single block structure. Here, ultrafine means that the prior austenite grain size is 2.0 μm or less. A single block structure means a structure in which only one block exists from the austenite grain. Adjacent blocks have different austenite parent phases. This can be confirmed by variant analysis using the results of EBSD (electron backscatter diffraction) observation of the structure.

Next, we will explain the manufacturing method of the ultrafine ferrite-cementite structure steel and ultrafine ferrite-austenite structure steel of the present invention.

In the manufacturing method of the ultrafine ferrite-cementite structure steel of the present invention, first, Fe powder, C powder, Si powder and Mn powder are adjusted to 0.05 wt%≦C≦0.3 wt%, 0.5≦Si≦2 wt%, 3 wt%≦Mn≦10 wt%, respectively, with the remainder being Fe, to obtain a raw material. Then, the raw material with the above chemical components is subjected to warm rolling processing at a rolling reduction rate of 50% or more in the temperature range of 300 to 600°C to obtain a fine ferrite-cementite structure, and Mn is concentrated in the cementite.

The manufacturing method of the ultrafine ferrite-austenite structure steel of the present invention is characterized in that the ultrafine ferrite-cementite structure steel manufactured by the above method is held at a temperature range of 625 to 750°C for 180 to 1200 seconds, and then water-cooled or air-cooled.

Warm rolling may be any of the following industrially performed methods: flat roll rolling in a thick steel plate production line, forging in an extra-thick steel plate production line, grooved roll rolling in a bar or steel wire production line, and shaped roll rolling in a bar or shaped steel line. Plastic strain is applied to the material by any of these processing methods. This warm rolling is performed at a temperature range of 300 to 600°C with a reduction ratio of 50 to 90%, resulting in an ultrafine ferrite-cementite structure. If the warm rolling temperature is less than 300°C, cementite is not formed, while if it exceeds 600°C, austenite is reverse transformed and formed. If the reduction ratio is less than 50%, the initial structure is work-hardened, and fine ferrite and fine cementite are not sufficiently developed. During this warm rolling, the martensite structure before rolling changes to a ferrite-cementite structure. The ferrite is ultrafine, and Mn is concentrated in the cementite.

The ultrafine ferrite-austenite structure steel obtained by warm rolling in this way is used as a pre-annealing structure for dual-phase region annealing, then held at a temperature range of 625-800°C for 1 second to 1 hour, and then water-cooled or air-cooled to produce a high-strength, high-ductility steel material with a single variant equiaxed martensite structure or an ultrafine martensite-austenite structure. If the reheating temperature is less than 625°C, the ferrite-cementite structure does not reverse transform completely, and a sufficient austenite fraction cannot be obtained. If it exceeds 800°C, the reverse-transformed austenite structure becomes coarse, and fine martensite cannot be obtained.

The manufacturing method of the ultrafine martensite structure steel and martensite-austenite structure steel of the present invention is characterized in that the ultrafine ferrite-cementite structure steel is held at a temperature range of 625 to 800°C for 1 to 600 seconds, and then water-cooled or air-cooled to produce a structure that is equiaxed martensite with a single block structure or an ultrafine equiaxed martensite-austenite structure steel with a single block structure.

The present invention will now be described more specifically with reference to examples.
<Example>
(1) Test material and structural observation Fe powder, C powder, Si powder, and Mn powder (0.1 wt%C-2 wt%Si-5 wt%Mn) were prepared as the main raw materials for melting, with the remaining Fe powder. These main raw materials for melting were melted using a high-frequency vacuum induction heating furnace, and a steel ingot measuring 100 mm long x 100 mm wide x 300 mm high was cast to be used as the material for the high-strength, high-ductility steel material of the present invention. This material was hot forged at 1200°C and formed into a square bar measuring 38 mm x 38 mm.

The material was then heated at 550°C for 1 hour, and reheated every 3 passes. It was then warm grooved rolled at a reduction rate of 90% until the cross section was 14 mm x 14 mm, and the structure was preliminarily made into ultrafine ferrite-cementite. Figure 1(a) shows a photograph of the obtained ultrafine ferrite-cementite structure taken with a field emission scanning electron microscope (FE-SEM). The ferrite grain size of the obtained structure was approximately 0.5 μm.

To make the structure of this rolled material martensite (comparative example), it was heated and held at 1200°C for 1 hour to austenitize it, and then air-cooled. Figure 1(b) shows an FE-SEM image of the austenite structure after air-cooling. In the following, ferrite is abbreviated as α, austenite as γ, cementite as θ, and martensite as M. The fine α+θ structure material and the comparative M structure material were heat-treated for 10 minutes, 30 minutes, and 1 hour, respectively, at 625°C and 675°C, which are in the α+γ two-phase region.

The solute C concentration in the γ grains was determined from the results of X-ray diffraction using a SmartLab made by Rigaku Corporation. The lattice constant can be calculated from the Bragg equation and the interplanar spacing d using formula (1).

Since Cu-Kα radiation was used, λ = 1.542 (A) was substituted, and θ was the angle obtained by measurement. Additionally, h, k, and l are plane indices (hkl). Using this formula, the lattice constants of γ in the γ phase (111), (200), and (220) were calculated. This was taken as the lattice constant a _γ (A) of γ using the method of extrapolating θ to 90 using (cos ² θ/sinθ) + (cos ² θ/θ), and was calculated using the Dyson and Holmes formula shown in formula (2).

Mn _γ was calculated by substituting the Mn concentration in γ measured by a field emission electron probe microanalyzer (FE-EPMA).

(2) Tensile Test The tensile test was performed using a round bar tensile test piece with a diameter of 3.5 mm and a parallel part length of 24.5 mm taken from the heat-treated square bar at a crosshead speed of 0.5 mm/min (strain rate of 4.9×10 ⁻⁴ s ⁻¹ ). Displacement was measured using an extensometer with a gauge length of 17.0 mm.

(3) In-situ transmission X-ray diffraction experiment for tensile test In order to directly analyze the effect of γ in the α+γ structure obtained by two-phase annealing on strength and ductility, the change in the γ volume fraction during the tensile test was measured. Tensile test specimens were taken from the heat-treated square bar so that the RD direction was the tensile direction. Plate-shaped tensile test specimens with a parallel section length of 12 mm and width of 2.5 mm were prepared. In order to transmit synchrotron radiation, the plate thickness was thinned to 0.5 mm. In order to measure the change in the γ volume fraction while proceeding with the tensile deformation in-situ, it is necessary to perform the diffraction profile change over a wide diffraction angle range in a short time. Therefore, in this example, the beamlines BL19B2 and BL15XU of the SPring-8 synchrotron radiation facility, which can use high-brightness X-rays, were used. The X-ray energy was 30 keV, and the beam size was 2.5 mm in the test piece width direction and 200 μm in the tensile direction. A small tensile tester was installed on a goniometer, and X-rays were incident from the normal direction of the tensile test specimen. A one-dimensional detector MYTHEN or an imaging plate (IP) or a two-dimensional detector PILATUS was installed behind the test piece, and an in-situ transmission X-ray diffraction experiment was performed at 30 keV. The crosshead speed of the tensile test was 0.245 mm/min, and the time resolution was 2 s.

The initial γ volume fraction was calculated using formula (3) from the theoretical hkl diffraction intensity and the peak area intensity ratio of the α phase (110), (200), (211) and the γ phase (111), (200), (220) at the obtained scattering angle. The subscript j indicates the jth diffraction peak used in the calculation, and n indicates the number of diffraction peaks of γ and α used in the calculation (n = 3).

The γ volume fraction during the tensile test was calculated from the product of the reduction rate from the initial peak area intensity sum of the (200), (211), and (311) peaks of the γ phase and the initial γ volume fraction.
Comparative Example
Conventional TRIP steel was prepared as a comparative material. That is, a cold-rolled steel sheet having a composition of 0.13%C-2%SI-1.6%Mn was held at 830°C for 60 seconds, slowly cooled to 700°C at 10K/s, and further quenched to 400°C at 60K/s, and then austempered by holding at 400°C for 10s, 60s, and 480s. In other respects, various tests were performed in the same manner as in the examples.
<Results>
Figure 1 (a) shows SEM images of the fine α+θ structure, which is the structure before dual-phase annealing, and (b) the M structure (comparative example). The fine α+θ structure is a structure consisting of fine elongated α grains and granular θ, and the white particles θ are dispersed in large quantities at the α grain boundaries and within the grains. In contrast, the M structure is composed of packets and blocks, and the blocks are plate-shaped grains, with a large number of block boundaries and no θ. Dual-phase annealing was performed on these fine α+θ and M structures.

Figure 2 shows the relationship between the two-phase annealing time and the structure. The annealing times are limited to 600s and 3600s, and the EBSD-IPF (inverse pole figure) map, grain boundary map, and phase map are shown. The ND direction of the observation surface is mapped, and the relationship between color and orientation is shown in a stereo triangle. The red lines on the grain boundary map are high-angle grain boundaries with a misorientation angle of 15° or more, blue are medium-angle grain boundaries with a misorientation angle of 5 to 15°, and light blue are low-angle grain boundaries with a misorientation angle of less than 5°. It is composed almost entirely of high-angle grain boundaries. The red and green colors on the phase map represent the ferrite phase (α) and the austenite phase (γ), respectively.

When the pre-annealing fine α+θ structure was annealed at 675°C in the two-phase region, it became a structure composed of α and γ, both of which were equiaxed, with the α and γ grain sizes being 2 μm or less, resulting in an ultrafine equiaxed grain structure. This is because γ was generated using the θ grains dispersed in the pre-annealing structure as nuclei, and the shape of the γ was spherical near the grain boundary triple junction and plate-like on the grain boundaries.

On the other hand, when the comparative example pre-annealed structure M was annealed in the two-phase region at 675°C, both α and γ had plate-like and needle-like structures, just like the pre-annealed structure. Martensite has a packet-block-lath structure, with blocks consisting of a group of laths with roughly the same crystal orientation, and blocks in the same packet share the (111) plane of the prior γ. The results of the phase map and IPF map obtained by EBSD analysis confirmed that acicular γ had formed at the block boundaries. Acicular γ had also formed within the blocks. This is thought to be the result of γ also forming at the lath boundaries of M. From this, it is thought that the nucleation sites of γ are the lath and block boundaries of M.

When the pre-annealing structure was fine α+θ, the α grains became slightly coarser as the annealing time increased from 600 s to 3600 s. On the other hand, the grain size and volume fraction of the γ grains increased. When the pre-annealing structure was M, it was confirmed that both the α grains and the γ grains became coarser.

Figure 3 shows the change in gamma volume fraction with changes in temperature and time in the two-phase annealing region. At all annealing temperatures, the gamma volume fraction increased as the annealing time increased. In addition, when comparing at the same holding time, the gamma volume fraction increased at 675°C. This is thought to be because the phase ratio of gamma in the α-γ equilibrium fraction increases at higher annealing temperatures and because diffusion also becomes faster. When the pre-annealing structure was fine α+θ, a large amount of gamma was obtained with annealing at 675°C for 3,600 seconds. In contrast, when the pre-annealing structure was M, the figure was 24.3%. Furthermore, when the pre-annealing structure was fine α+θ, the gamma volume fraction was high at 20% even with a short annealing time of 600 seconds, which is twice the gamma volume fraction obtained when the pre-annealing structure was M, at 10%.

Figure 4 is a schematic diagram of the process of γ generation by reverse transformation. As shown in Figure 4 (b), in the case of reverse transformation from the pre-annealing structure M, a supersaturated C solid solution, θ first precipitates in large numbers at lath and block boundaries from M, which is a supersaturated C solid solution, and these combine and connect to grow into fine needle-like θ, and γ nucleates from the needle-like θ. As a result, acicular γ is thought to be generated from the needle-like θ. On the other hand, as shown in Figure 4 (a), in the case of reverse transformation from the pre-annealing structure of fine α + θ, γ nucleates from θ that is originally dispersed in large amounts at the α grain boundaries and within the grains, and equiaxed γ is thought to be generated. In this way, it is believed that the finely dispersed θ acts effectively as a preferential nucleation site for reverse transformation γ, making it possible to form a large amount of γ (21%) in an extremely short time (600 s). For this reason, it is considered that it is extremely important in the present invention that the structure before two-phase annealing is ferrite-cementite.

According to Thermo-Calc calculations, the composition of θ at 550°C, which is the rolling temperature when producing the fine α+θ structure before annealing, is ( _Fe0.49 , _Mn0.51 ) _3C , and it is believed that Mn is highly concentrated in θ. On the other hand, when the structure before annealing is M, it is produced by quenching from 1200°C, so it is believed that Mn and C are uniformly present in the martensite.

Figures 5 and 6 show SEM and TEM images of the fine α+θ structure before annealing, and the α+θ structure and M structure before annealing, which were heated to 675°C and immediately cooled in water without holding. Samples for TEM observation were prepared using the extraction replica method. As can be seen from Figure 5(a), in the case of the fine α+θ structure, the elongated α has recrystallized and become equiaxed. θ can also be confirmed. In Figure 5(b), it can be seen that many small spherical θ have precipitated, and precipitation of an acicular structure is also observed.

As is clear from Figure 6 (a), a large amount of θ grains are present in the fine α + θ structure before annealing. Figure 6 (b) shows that the number of θ grains has decreased and their size has also become smaller immediately after heating to 675°C. This is because the formation of γ has begun due to reverse transformation. When the structure before annealing is M, the acicular structure observed in Figure 5 (b) is thought to be θ from Figure 6 (c). Some of it is also thought to be γ. From the results of the TEM-EDS analysis in Figure 7, θ in the fine α + θ structure before annealing contained 50.8 (wt%) Mn, which is almost the stoichiometric composition. However, the Mn concentration in θ decreased immediately after heating to the two-phase region, which is thought to be because Mn had already migrated to the reverse transformation γ. When the structure before annealing is M, the acicular θ was formed during the heating process, and the Mn concentration in θ was 36.6 (wt%), which was lower than in the case of fine α + θ.

Therefore, when the pre-annealing structure is fine α + θ, Mn is already present in high concentrations in the finely dispersed θ, and it is believed that this high concentration of Mn enables the formation of highly stable γ in a short period of time. Furthermore, the concentration of Mn in this cementite is a phenomenon that is essential for the rapid formation of austenite.

Figure 8 shows the results of Mn concentration mapping by FE-EPMA for the α+γ structure after dual-phase annealing. Figure 8(a) shows the results when the structure before annealing was fine α+θ and annealed at 675°C for 3600 seconds, confirming that the Mn concentration in γ was extremely high. Quantitative analysis showed that the Mn concentration was 9.6 wt% on average at the five points that showed high concentrations.

On the other hand, Figure 8 (b) shows the results when the pre-annealing structure was M and annealing was performed at 675°C for 600 seconds. The Mn concentration mapping confirmed that the Mn concentration was relatively uniform.

Figure 9 shows the nominal stress-nominal strain curves obtained in the round bar tensile tests of the four α+γ structures shown in Figure 2. When the pre-annealing structure was fine α+θ, discontinuous yielding occurred, whereas when the pre-annealing structure was M, continuous yielding occurred, and the shapes of the nominal stress-nominal strain curves were different.

On the other hand, in both pre-annealing structures, as the two-phase annealing time was increased, the yield stress decreased and the tensile strength increased. This is because, as is clear from Figure 2, when the pre-annealing structure is α+θ, the α grains become coarse as the annealing time is increased, and it is thought that the α grain size determines the yield point. When the pre-annealing structure is M, the blocks also clearly become coarse, and it is thought that this influence caused the yield point to decrease with long-term annealing. On the other hand, the tensile strength is higher with longer annealing times because the γ volume fraction increases, as is clear from Figure 3.

When the pre-annealing structure is α+θ, annealing for a very short time (600 s) showed high yield stress (870 MPa), tensile strength (960 MPa) and total elongation (33.5%), with a strength-ductility balance of 32,200 MPa%, and excellent mechanical properties exceeding 30,000 MPa%, which is an indicator of excellent strength-ductility balance. Furthermore, even with an annealing time of 3,600 s, although the yield stress (770 MPa) decreased slightly, high tensile strength (1,090 MPa) and total elongation (32.4%) were shown, and the strength-ductility balance was 35,300 MPa, showing excellent mechanical properties. On the other hand, in the comparative example, where the pre-annealing structure was M, an annealing time of 600 s showed yield stress (600 MPa), tensile strength (897 MPa), and total elongation (25.4%), resulting in a strength-ductility balance of 22,780 MPa%, while an annealing time of 3,600 s showed yield stress (550 MPa), tensile strength (970 MPa), and total elongation (32.3%), resulting in a strength-ductility balance of 31,330 MPa%.

The α+θ pre-annealing structure provides a better balance of strength and ductility than M. As is clear from the results in Figure 3, the main reason for this is thought to be that the γ volume fraction is higher when the pre-annealing structure is α+θ than when it is M.

Figure 10 shows the nominal stress-nominal strain curves of plate-shaped tensile test pieces of α+γ structured steel with α+θ and M pre-annealing conditions of 675°C x 3600s, and the change in γ volume fraction during tensile deformation. For comparison, the nominal stress-nominal strain curves of conventional TRIP steel (austempering time 480s) and the change in γ volume fraction during tensile testing are also shown. As with the round bar tensile test, when the pre-annealing structure was fine α+θ, the strength and ductility were superior to when the pre-annealing structure was M. Although the processing-induced transformation progresses rapidly from the start of plastic deformation, when the pre-annealing structure was α+θ, γ remained up to the higher strain range, which is thought to be the mechanism for high strength and high ductility. Conventional TRIP steel has a low γ volume fraction of 10%, but γ remained up to the higher strain range, demonstrating the TRIP effect.

The change in austenite stability caused by the two-phase annealing conditions is shown in Table 1.

As shown in Table 1, the solute carbon concentration in γ (C _γ ) decreases as the annealing time increases. This is thought to be because the γ volume fraction (γ _R ) increases with the annealing time. The product of the solute carbon concentration and the γ volume fraction, γ _R ×C _γ , is an index that indicates how much of the 0.1% added C is dissolved in γ, and the maximum is 0.1. When the pre-annealing structure is M, γ _R ×C _γ is 0.033 to 0.057, which is thought to be a state in which a large amount of solute C remains in M. On the other hand, when the pre-annealing structure is fine α + θ, γ _R ×C _γ is 0.073 to 0.091, which means that θ has almost changed to γ at an annealing time of 3600 s. It can be said that θ has efficiently changed to almost γ even at an annealing time of 600 s. When the prior structure is α+θ, the added C is precipitated as θ in advance, which is thought to be the mechanism by which C can be efficiently dissolved in a large amount of γ. The amount of C added in the conventional TRIP steel is 0.13%, while γ _R ×C _γ is 0.097, and the C concentration ratio in γ is 74%, which is approximately the same as 73% when the prior structure is α+θ and the annealing time is 600 s. This means that making the prior structure α+θ gives an efficiency comparable to that of austempering.

The stability of gamma is considered using the Md30 formula shown in formula (4). Md30 is the temperature at which 50% of the entire structure transforms into the M phase when a true tensile strain of 0.30 is applied to a sample with a single-phase gamma structure; the higher this value, the more unstable gamma is.

When the pre-annealing structure was α+θ and the annealing conditions were 675°C x 3600s, and when the pre-annealing structure was M and the annealing conditions were 675°C x 3600s, the Md30 was 200°C and 227°C, respectively. In the case of the pre-annealing structure α+θ, although the gamma volume fraction was high, the stability was also high, so it is believed that gamma remained up to the high strain region compared to when the pre-annealing structure was M. This is believed to have resulted in high ductility. The Md30 of conventional TRIP steel was -76°C. Although the gamma volume fraction was low and the tensile strength was low, it is believed that the extremely high stability resulted in high ductility.

In Fig. 11, the relationship between the strength and ductility balance of the present α+γ structure steel and the conventional TRIP steel is plotted based on the results of the plate tensile test shown in Fig. 10. The strength and ductility balance of the present α+γ steel with a 0.1%C-2%Si-5%Mn composition is superior to that of the conventional TRIP steel. The conventional TRIP steel exhibits high ductility by concentrating extremely high C in a small amount of γ through austempering, but has low strength due to its low γ volume fraction. In contrast, the α+γ steel with a 0.1%C-2%Si-5%Mn composition has a fine α+θ structure before annealing, which gives it a higher γ volume fraction than the conventional TRIP steel. Although the amount of solute C in γ is low, the effect of solute Mn stabilizes γ, resulting in high ductility. Note that the total elongation depends on the size of the test piece, so this comparison was made using the same plate tensile test piece. Even with an extremely short annealing time of 600 s, it showed very excellent mechanical properties.
<Strength test of ultrafine ferrite-austenite structure steels depending on carbon content>
2%Si-5%Mn steels with varying C contents of 0.075, 0.1, 0.15, 0.2, and 0.3 wt% were prepared by vacuum melting and hot forging, and ultrafine ferrite (α) + austenite (γ) structures were prepared by warm groove rolling. These samples were annealed in the two-phase region at 675°C and 700°C for 1 hour each to prepare test materials. Strength and ductility were examined by round bar tensile tests, and structures were observed by SEM-EBSD. Furthermore, thin plate test pieces (GL = 12 mm, T = 0.4 mm) were prepared, and in-situ XRD was performed during tensile deformation using high brightness X-rays from SPring-8. The γ volume fraction during the tensile test (CHS = 0.245 mm/min) was calculated from the integrated intensity ratio of each peak.

Figures 12 and 13 show the nominal stress-nominal strain curves (solid lines) and the gamma volume fraction (dashed lines) during tensile testing for test materials annealed at 675°C and 700°C, respectively. Comparing the physical properties of the 0.15C test materials from Figures 12 and 13, the 675°C annealed material showed a tensile strength of 1000 MPa and a total elongation of 40%, with 5% gamma remaining at fracture, whereas the 700°C annealed material showed a tensile strength of 1480 MPa and a total elongation of 30%.

FIG. 14 shows the relationship between γ volume fraction and tensile strength. It was confirmed from FIG. 14 that the tensile strength increases as the γ volume fraction increases. FIG. 15 shows the relationship between V γ xCin _γ , which is the Cin _γ product of the γ volume fraction and the solute carbon concentration in γ of the 675° C. annealed material, and the uniform elongation. As _{V γ xCin γ} increases, the uniform elongation increases. Focusing on the circled area in FIG. 14, even though the solute carbon concentration in γ is the same, there is a difference of about 10% in the uniform elongation, and in order to obtain a large uniform elongation of about 40%, it is necessary to increase V _γ xCin _γ . For example, V _γ xCin _γ ≧0.15 is preferably exemplified. From these, it was found that the tensile strength is correlated with the γ volume fraction, and the uniform elongation is correlated with V _γ xCin _γ . In addition, as shown in FIG. 16, the 700°C annealed material containing 0.15% C achieved a large γ volume fraction of 45% and V _γ xCin _γ of 0.15%, thereby exhibiting excellent ultrahigh strength and high ductility.

Claims (3)

An ultrafine equiaxed grain ferrite-austenite structure steel having a C, Si and Mn content of 0.05 wt%≦C≦0.3 wt%, 0.5 wt%≦Si≦2 wt%, 3 wt%≦Mn≦10 wt%, with the balance being Fe and unavoidable impurities , a ferrite grain size of 2.0 μm or less, and an austenite grain size of 2.0 μm or less ,
An ultrafine ferrite-austenite structure steel, characterized in that an austenite volume fraction is 40% or more, VγxCinγ≧0.07, where Vγ represents the austenite volume fraction, Cinγ represents a solute carbon concentration in austenite , and the steel has a tensile strength of 1400 MPa or more and a total elongation of 25% or more. A raw material having a martensite structure in which the contents of C, Si and Mn are 0.05 wt%≦C≦0.3 wt%, 0.5 wt%≦Si≦2 wt%, 3 wt%≦Mn≦10 wt%, respectively, with the balance being Fe and unavoidable impurities, is heated to a temperature range of 300 to 600° C. and subjected to warm rolling at a rolling reduction rate of 50% or more, thereby forming a fine ferrite-cementite structure with a ferrite grain size of 2 μm or less , concentrating Mn in the cementite , and obtaining an ultrafine ferrite-cementite structure steel in which the value of x in the chemical formula of the cementite (Fe1-x, Mnx)3C is 0.3 or more and less than 1 ;
The ultrafine ferrite-cementite structure steel is held at a temperature range of 625 to 750° C. for 180 to 1200 seconds, and then water-cooled or air-cooled;
A method for producing an ultrafine ferrite-austenite structure steel, comprising: A material having C, Si and Mn contents of 0.05 wt%≦C≦0.3 wt%, 0.5 wt%≦Si≦2 wt%, 3 wt%≦Mn≦10 wt%, respectively, with the balance being Fe and unavoidable impurities, is heated to a temperature range of 300 to 600° C. and subjected to warm rolling processing at a rolling reduction rate of 50% or more to form a fine ferrite-cementite structure, and Mn is concentrated in the cementite to obtain an ultrafine ferrite-cementite structure steel;
The ultrafine ferrite-cementite structure steel is held in a two-phase temperature range of 625 to 800°C for 1 to 600 seconds, and then water-cooled or air-cooled to produce a structure that is an ultrafine equiaxed martensite-austenite structure steel having a single block structure;
A method for producing martensite -austenite steel, comprising:

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