High Strength Dual-Phase TRIP Steel and Method for Making Same

Abstract

A high strength dual-phase TRIP steel includes 8-12 wt. % Mn, 0.4-0.6 wt. % C, 1-3 wt. % Al, 0.5-1 wt. % V, and a balance of Fe. A method for making the high strength dual-phase TRIP steel is also provided.

Description

BACKGROUND

1. Technical Field

The subject matter herein generally relates to a high strength dual-phase TRIP steel, and a method for making the high strength dual-phase TRIP steel.

2. Description of Related Art

Automotive industry is one of the leading greenhouse gas emitters globally. One effective way for the automotive industry to reduce greenhouse gas emission is by reducing vehicles' weight, which can be achieved by using lightweight materials. The most common lightweight materials currently used in the automotive industry are advanced high strength steels (AHSS). The high strength offered by AHSS leads to thinner components and lighter cars can be made. Nevertheless, AHSS should also possess a certain amount of ductility, so that complex automotive parts can be formed by stamping technology at room temperature. Therefore, both the automotive industry and the steel industry are currently searching for new AHSS possessing simultaneously high strength and good ductility.

As can be seen in FIG. 1, the first generation of AHSS has been referred to a variety of steels including dual phase (DP) steel, transformation-induced plasticity (TRIP) steel, complex-phase (CP) steel, and martensitic (MART) steel. The second generation AHSS mainly includes twinning-induced plasticity (TWIP) steel with high manganese content. Recent global interest is to develop the “Third Generation” of AHSS, i.e. steels with strength-ductility combinations between the first and second generation as shown FIG. 1.

The Medium Mn steels which contain the Mn content from 5 to 12 wt. % are considered to be the potential candidates to achieve the targeted mechanical properties of third generation of AHSS. Shi et al. reported that the 5 Mn steel (Fe-0.2C-5Mn, wt. %) which contains the retained austenite grains embedded in the ferrite matrix can have a tensile strength of 1420 MPa and a total elongation of 31%. It is suggested that both the amount and stability of retained austenite grains are vital for the work hardening behavior and therefore the mechanical properties. The stability of austenite grain is governed by chemical compositions, grain size, matrix strength, morphology and free surface. Lee et al. reported an exceptional elongation (−65%) of the 10 Mn steel (Fe-10Mn-0.3C-3Al-2Si, wt. %) after intercritical annealing. They ascribed the excellent tensile ductility to the sequence of twinning induced plasticity effect followed by the TRIP effect. However, the yield strength of the above 5 Mn steel and 10 Mn steel is low, which are 600 MPa and 800 MPa, respectively.

Therefore, a need exists in the industry to overcome the described problems.

SUMMARY

The disclosure is to offer a high strength dual-phase TRIP steel, and a method for making the high strength dual-phase TRIP steel.

A high strength dual-phase TRIP steel comprises 8-12 wt. % Mn, 0.4-0.6 wt. % C, 1-3 wt. % Al, 0.5-1 wt. % V, and a balance of Fe.

Preferably, the high strength dual-phase TRIP steel comprises 10 wt. % Mn, 0.47 wt. % C, 2 wt. % Al, 0.7 wt. % V, and a balance of Fe.

Preferably, the high strength dual-phase TRIP steel consists of ferrite and retained austenite phases.

Preferably, wherein the high strength dual-phase TRIP steel consists of two types of retained austenite and fine-grained ferrite.

Preferably, wherein one type of the retained austenite is fine-grained austenite with a grain size which is smaller than 1 μm, the other type is the large-grained austenite with a size which is larger than 10 μm.

Preferably, a volume fraction of austenite contained in the high strength dual-phase TRIP steel before a tensile test is 70-90%, a volume fraction of ferrite contained in the high strength dual-phase TRIP steel before the tensile test is 10-30%.

Preferably, the volume fraction of austenite contained in the high strength dual-phase TRIP steel before the tensile test is 80.8%, the volume fraction of ferrite contained in the high strength dual-phase TRIP steel before the tensile test is 19.8%.

Preferably, a volume fraction of the fine-grained austenite contained in the retained austenite is 60-80%, a volume fraction of the large-grained austenite contained in the retained austenite is 20-40%.

Preferably, after the high strength dual-phase TRIP steel is deformed, the volume fraction of austenite drops to 30-60%, the volume fraction of martensite increases to 30-60% and the volume fraction of the ferrite remains still.

Preferably, the volume fraction of austenite drops to 52.7%, the volume fraction of martensite increases to 28.1% and the volume fraction of the ferrite remains 19.8%.

Preferably, the volume fraction of the fine-grained austenite contained in the retained austenite is 67%, the volume fraction of the large-grained austenite contained in the retained austenite is 33%.

Preferably, the high strength dual-phase TRIP steel includes vanadium carbide precipitations with a size of about 20-30 nm.

A method for making the high strength dual-phase TRIP steel, comprising:

ingots are provided, the ingots comprises 8-12 wt. % Mn, 0.4-0.6 wt. % C, 1-3 wt. % Al, 0.5-1 wt. % V, and a balance of Fe;

the ingots are hot rolled to produce a plurality of thick steel sheets with a thickness of 5-6 mm, then the steel sheets are treated by an air cooling process;

the steel sheets are warm rolled at a temperature of about 150-200 with a thicknesses reduction of about 15-44%;

the steel sheets are cold rolled at a room temperature with the thicknesses reduction of about 7%;

the steel sheets are annealed at the temperature of about 625-660 for about 10-300 mins to form the high strength dual-phase TRIP steel.

Preferably, a starting hot rolling temperature is 1100-1300, and a finishing hot rolling temperature is 800-1000, the thickness of each steel sheet is 3-6 mm.

Preferably, the steel sheets are finally water quenched to the room temperature after the anneal process.

Preferably, the high strength dual-phase TRIP steel comprises 8-12 wt. % Mn, 0.4-0.6 wt. % C, 1-3 wt. % Al, 0.5-1 wt. % V, and a balance of Fe.

Preferably, the high strength dual-phase TRIP steel comprises 10 wt. % Mn, 0.47 wt. % C, 2 wt. % Al, 0.7 wt. % V, and a balance of Fe.

Preferably, the high strength dual-phase TRIP steel consists of ferrite and retained austenite phases, the retained austenite phase consists of fine-grained austenite with a grain size which is smaller than 1 μm, and large-grained austenite with a size which is larger than 10 μm.

Preferably, a volume fraction of austenite contained in the high strength dual-phase TRIP steel before a tensile test is 70-90%, a volume fraction of ferrite contained in the high strength dual-phase TRIP steel before the tensile test is 10-30%, a volume fraction of the fine-grained austenite contained in the retained austenite is 60-80%, a volume fraction of the large-grained austenite contained in the retained austenite is 20-40%.

Preferably, the high strength dual-phase TRIP steel includes vanadium carbide precipitations with a size of 20-30 nm.

Compared with the first and the second generation of AHSS, the invention has the following technical effects: the TRIP effect occurred in the high strength dual-phase TRIP steel can improve the strength of the high strength dual-phase TRIP steel. Furthermore, the V element and the C element can react with each other to form vanadium carbide precipitations, the vanadium carbide precipitations can also improve the strength of the steel.

BRIEF DESCRIPTION OF THE DRAWINGS

Many aspects of the present embodiments can be better understood with reference to the following drawings. The components in the drawings are not necessarily drawn to scale, the emphasis instead being placed upon clearly illustrating the principles of the present embodiments. Moreover, in the drawings, all the views are schematic, and like reference numerals designate corresponding parts throughout the several views.

FIG. 1 is an isometric view of different generations of automotive steels.

FIG. 2 is a flow chart of a method for making the high strength dual-phase TRIP steel according to an exemplary embodiment.

FIG. 3 is a schematic diagram for various thermomechanical treatments.

FIG. 4a shows tensile results of the high strength dual-phase TRIP steels according to a first exemplary embodiment.

FIG. 4b shows engineering stress-strain curves of the high strength dual-phase TRIP steels according to a second exemplary embodiment.

FIG. 4c shows the engineering stress-strain curves of the high strength dual-phase TRIP steels according to a third exemplary embodiment.

FIG. 4d shows the engineering stress-strain curves of the high strength dual-phase TRIP steels according to a fourth exemplary embodiment.

FIG. 5 presents XRD results of the high strength dual-phase TRIP steels with 0% strain, 7% strain and fracture.

FIG. 6 presents an evolution of the retained austenite volume fraction of the high strength dual-phase TRIP steel during a tensile test according to an exemplary embodiment.

FIGS. 7a and 7b are EBSD band contrast and phase images of an initial microstructure of the high strength dual-phase TRIP steel during the tensile test according to an exemplary embodiment.

FIG. 8 shows a microstructure of the fractured high strength dual-phase TRIP steel according to an exemplary embodiment.

DETAILED DESCRIPTION

The disclosure is illustrated by way of example and not by way of limitation in the figures of the accompanying drawings, in which like reference numerals indicate similar elements. It should be noted that references to “an” or “one” embodiment in this disclosure are not necessarily to the same embodiment, and such references can mean “at least one” embodiment.

According to an exemplary embodiment, a high strength dual-phase TRIP steel (transformation induced plasticity steel, TRIP) for automotive applications comprises, by weight percent: 8-12 wt. % Mn, 0.4-0.6 wt. % C, 1-3 wt. % Al, 0.5-1 wt. % V, and a balance of Fe.

In an exemplary embodiment, the high strength dual-phase TRIP steel comprises, by weight percent: 10 wt. % Mn, 0.47 wt. % C, 2 wt. % Al, 0.7 wt. % V, and a balance of Fe.

The high strength dual-phase TRIP steel consists of ferrite and retained austenite phases. The austenite phase contained in high strength dual-phase TRIP steel is stability, so that the TRIP effect can take place gradually in the retained austenite.

In at least one exemplary embodiment, the high strength dual-phase TRIP steel consists of two types of retained austenite and fine-grained ferrite. One type of the retained austenite is fine-grained austenite with a grain size that is smaller than 1 μm. The other type is the large-grained austenite with a size larger than 10 μm. A volume fraction of the fine-grained austenite contained in the retained austenite is 60-80%, a volume fraction of the large-grained austenite contained in the retained austenite is 20-40%.

In at least one exemplary embodiment, the volume fraction of the fine-grained austenite contained in the retained austenite is 67%, a volume fraction of the large-grained austenite contained in the retained austenite is 33%.

A volume fraction of the austenite contained in the high strength dual-phase TRIP steel before a tensile test is 70-90%, a volume fraction of ferrite contained in the high strength dual-phase TRIP steel before the tensile test is 10-30%, in at least one exemplary embodiment, the volume fraction of the austenite contained in the high strength dual-phase TRIP steel before a tensile test is 80.8%, the volume fraction of ferrite contained in the high strength dual-phase TRIP steel before the tensile test is 19.2%. After a deformation, the volume fraction of the austenite drops to 30-60%, that suggests an occurrence of the TRIP effect, which can also be observed from an EBSD result of a fracture structure of the high strength dual-phase TRIP steel. After the deformation, the volume fraction of martensite increases to 30-60% and the volume fraction of the ferrite remains still. The TRIP effect results in high working hardening rate, high ultimate tensile strength and good elongation. In at least one exemplary embodiment, after the deformation, the volume fraction of austenite drops to 52.7%, the volume fraction of martensite increases to 28.1% and the volume fraction of the ferrite remains 19.2%.

It is to be understood that, the TRIP effect occurred in the steel can improve the work hardening behavior of the steel. As a result, the strength of the steel can be increased without losing ductility. Furthermore, the formation of vanadium carbide precipitations during the annealing process can provide the precipitation hardening to strengthen the steel.

The vanadium carbide precipitations are nanosized, with a diameter of about 20-30 nm. Such a proper size of the precipitations can efficiently increase the strength of steel by Orowan bypassing mechanisms. The high strength dual-phase TRIP steel can have a high yield stress, high work hardening rate, high ultimate tensile stress and good ductility. Nanosized vanadium carbide precipitations contribute to the high yield strength of the high strength dual-phase TRIP steel.

It is to be understood that, the yield stress of the high strength dual-phase TRIP steel is about 1400 MPa, the ultimate tensile strength of the high strength dual-phase TRIP steel is about 1586 MPa, the total elongation of the high strength dual-phase TRIP steel is about 15.2%. It is noted that a uniform elongation of the high strength dual-phase TRIP steel is almost the same as the total elongation of the high strength dual-phase TRIP steel.

Referring to FIGS. 2-3, the invention discloses a thermomechanical method for making the high strength dual-phase TRIP steel. The method is provided by way of example, as there are a variety of ways to carry out the method. The method described below can be carried out using the configurations illustrated in FIG. 2, for example, and various elements of the figures are referenced in explaining method. Each block shown in FIG. 2 represents one or more process, methods or subroutines, carried out in the method. Furthermore, the order of blocks is illustrative only and the blocks can change according to the present disclosure. Additional blocks can be added or fewer blocks can be utilized, without departing from this disclosure. The method for making the hinge can begin at block 201.

At block 201, ingots are provided. It is to be understood that, the ingots comprises, by weight: 8-12 wt. % Mn, 0.4-0.6 wt. % C, 1-3 wt. % Al, 0.5-1 wt. % V, and a balance of Fe.

At block 202, the ingots are hot rolled to produce a plurality of 3-6 mm thick steel sheets, followed by an air cooling process. It is to be understood that, a starting hot rolling temperature is 1100-1300, and a finishing hot rolling temperature is 800-1000. In at least one exemplary embodiment, the thickness of each steel sheet is 4 mm, the starting hot rolling temperature is 1200, and the finishing hot rolling temperature is 900.

At block 203, the steel sheets are warm rolled at a temperature of about 150-200 with a thicknesses reduction of about 15-44%. The warm rolling process can minimize the transformation of austenite to martensite, and can be employed to avoid the occurrence of cracks during the warm rolling process.

At block 204, the steel sheets are cold rolled at the room temperature with a thicknesses reduction of about 7%.

At block 205, the steel sheets are then annealed at a temperature of about 625-660 for about 10-300 min. The vanadium carbide precipitations are formed during the annealing process.

After the annealing process, there are large amount of the retained austenite grains which are metastable in steel. During tensile deformation, these retained austenite can transform to martensite by the stress assisted or strain assisted process. The corresponding martensitic transformation provides the TRIP effect, which results in a high working hardening rate, a high ultimate tensile strength and a good elongation.

At block 206, the steel sheets are finally water quenched to the room temperature.

It is to be understood that, after the steel sheets are cold rolled, the steel sheets can be wire-cut from the rolled sheets with the tensile axis aligned parallel to the rolling direction to achieve a plurality of tensile test samples, the tensile test samples can be tested through a standard test Method for tension testing of metallic materials (ASTM-E8).

FIG. 4a shows tensile results of the high strength dual-phase TRIP steels according to a first exemplary embodiment. In detail, four test samples according to the first exemplary embodiment are made by the following steps: the four test sheets are warm rolled at the temperature of about 150 with the thicknesses reduction of about 44%, then the test sheets are cold rolled at the room temperature with the thicknesses reduction of about 7%, finally the steel sheets are respectively annealed at the temperature of about 660 for about 30 min (referring to curve (a) of FIG. 4a), at the temperature of about 660 for about 10 min (referring to curve (b) of FIG. 4a), at the temperature of about 625 for about 300 min (referring to curve (c) of FIG. 4a), or at the temperature of about 625 for about 60 min (referring to curve (d) of FIG. 4a). It seems that the annealing at the 625 for about 300 min provides a promising combination of high tensile strength and good ductility of TRIP steel. FIG. 4b shows engineering stress-strain curves of the high strength dual-phase TRIP steels according to a second exemplary embodiment. In detail, four test samples according to the second exemplary embodiment are made by the following steps: the four test sheets are warm rolled at the temperature of about 150 with the thicknesses reduction of about 15%, then the test sheets are cold rolled at the room temperature with the thicknesses reduction of about 7%, finally the steel sheets are respectively annealed at the temperature of about 660 for about 30 min (referring to curve (a) of FIG. 4b), at the temperature of about 660 for about 10 min (referring to curve (b) of FIG. 4b), at the temperature of about 625 for about 300 min (referring to curve (c) of FIG. 4b), or at the temperature of about 625 for about 60 min (referring to curve (d) of FIG. 4b). Again, it seems that the annealing at the 625 for about 300 min provides the best tensile properties.

FIG. 4c shows the engineering stress-strain curves of the high strength dual-phase TRIP steels according to a third exemplary embodiment. In detail, four test samples according to the third exemplary embodiment are made by the following steps: the four test sheets are warm rolled at the temperature of about 200 with the thicknesses reduction of about 30%, then the test sheets are cold rolled at the room temperature with the thicknesses reduction of about 7%, finally the steel sheets are respectively annealed at the temperature of about 660 for about 30 min (referring to curve (a) of FIG. 4c), at the temperature of about 660 for about 10 min (referring to curve (b) of FIG. 4c), at the temperature of about 625 for about 300 min (referring to curve (c) of FIG. 4c), or at the temperature of about 625 for about 60 min (referring to curve (d) of FIG. 4c). It can be seen that the annealing at temperature of 625 for about 300 min is beneficial for the tensile properties.

FIG. 4d shows the engineering stress-strain curves of the high strength dual-phase TRIP steels according to a fourth exemplary embodiment. In detail, four test samples according to the fourth exemplary embodiment are made by the following steps: the four test sheets are warm rolled at the temperature of about 200 with the thicknesses reduction of about 15%, then the test sheets are cold rolled at the room temperature with the thicknesses reduction of about 7%, finally the steel sheets are respectively annealed at the temperature of about 660 for about 30 min (referring to curve (a) of FIG. 4d), at the temperature of about 660 for about 10 min (referring to curve (b) of FIG. 4d), at the temperature of about 625 for about 300 min (referring to curve (c) of FIG. 4d), or at the temperature of about 625 for about 60 min (referring to curve (d) of FIG. 4d). Although the tensile properties are more or less stable at these two annealing temperatures, the corresponding tensile properties are not as good as the one in curve (c) in FIG. 4a.

FIG. 5 presents XRD results of the high strength dual-phase TRIP steels with 0% strain (referring to curve (a) of FIG. 5), 7% strain (referring to curve (b) of FIG. 5) and fracture (referring to curve (c) of FIG. 5). The austenite γ(111) and γ(220) peaks dramatically decrease and the peak γ(200) even disappears after 7% strain. But there is no obvious difference between the samples subjected to 7% strain and fracture. The volume fraction of retained austenite can be calculated from the XRD results based on the American Society for Testing and Materials (ASTM) E975-03.

An evolution of the retained austenite volume fraction of the high strength dual-phase TRIP steel during the tensile test is shown in FIG. 6, wherein the corresponding engineering stress-strain curve is included. The initial retained austenite volume fraction is about 70-90%, the initial ferrite volume fraction is about 10-30%. The decrease of the austenite volume fraction after 7% strain suggests the occurrence of TRIP effect after the yielding and contributes to the high work hardening rate during the tensile test. However, the retained austenite volume fraction almost remained as a constant from 7% strain up to fracture. The slope of the tensile curve starts to increase after about 9% strain, indicating that other strengthening mechanism shall contribute to sustain the high work hardening rate, which could be twinning-induced plasticity (TWIP) effect. In other words, nanotwins may be generated after 9% straining. The nanotwin boundaries can act as barriers to dislocation glide, leading to the pile-up of dislocations at the twin boundaries and to the dynamic Hall-Petch effect. Therefore, the formation of nanotwin during tensile tests significantly enhances the work hardening rate, resulting in excellent uniform elongation. To investigate the mechanical property, uni-axial tensile tests are carried out at room temperature with an initial strain rate of about 10⁻³ s⁻¹. Interrupted tensile tests were applied to the high strength dual-phase TRIP steel at different engineering strains including 0% strain, 7% strain and fracture strain. For microstructure observation, an electron back-scattering diffraction (EBSD) measurement is performed in FEG SEM LEO 1530 at 20 kV. The data was processed by the HKL Channel 5. For the phase identification, the X-Ray diffraction (XRD) using Cu K_α radiation with wavelength 1.5405(6) Å is performed. The high strength dual-phase TRIP steels used for EBSD and XRD analysis are electro-polished after the mechanical finish of 1 m in a solution of 25% perchloric acid and 75% ethanol (vol. %) at room temperature.

FIGS. 7a and 7b are the EBSD band contrast and phase images of the initial microstructure of the high strength dual-phase TRIP steel. FIGS. 7a and 7b show that the initial microstructure of the high strength dual-phase TRIP steel consists of two types of retained austenite and fine-grained ferrite. One type of the retained austenite is fine-grained austenite with a grain size that is smaller than 1 μm with a volume fraction ranging 60-80%. The other type is the large-grained austenite with a size larger than 10 μm with a volume fraction ranging 20-40%.

FIG. 8 shows the microstructure of the fractured high strength dual-phase TRIP steel. A large amount of austenite is remained after the tensile test, which is consistent with the XRD results.

Although the features and elements of the present disclosure are described as embodiments in particular combinations, each feature or element can be used alone or in other various combinations within the principles of the present disclosure to the full extent indicated by the broad general meaning of the terms in which the appended claims are expressed.

Claims

1. A high strength dual-phase TRIP steel, comprising: 8-12 wt. % Mn, 0.4-0.6 wt. % C, 1-3 wt. % Al, 0.5-1 wt. % V, and a balance of Fe.

2. The high strength dual-phase TRIP steel of claim 1, wherein the high strength dual-phase TRIP steel comprises 10 wt. % Mn, 0.47 wt. % C, 2 wt. % Al, 0.7 wt. % V, and a balance of Fe.

3. The high strength dual-phase TRIP steel of claim 1, wherein the high strength dual-phase TRIP steel consists of ferrite and retained austenite phases.

4. The high strength dual-phase TRIP steel of claim 3, wherein the high strength dual-phase TRIP steel consists of two types of retained austenite and fine-grained ferrite.

5. The high strength dual-phase TRIP steel of claim 4, wherein one type of the retained austenite is fine-grained austenite with a grain size that is smaller than 1 μm, the other type is the large-grained austenite with a size larger than 10 μm.

6. The high strength dual-phase TRIP steel of claim 3, wherein a volume fraction of austenite contained in the high strength dual-phase TRIP steel before a tensile test is 70-90%, a volume fraction of ferrite contained in the high strength dual-phase TRIP steel before the tensile test is 10-30%.

7. The high strength dual-phase TRIP steel of claim 6, wherein the volume fraction of austenite contained in the high strength dual-phase TRIP steel before the tensile test is 80.8%, the volume fraction of ferrite contained in the high strength dual-phase TRIP steel before the tensile test is 19.2%.

8. The high strength dual-phase TRIP steel of claim 6, wherein a volume fraction of the fine-grained austenite contained in the retained austenite is 60-80%, a volume fraction of the large-grained austenite contained in the retained austenite is 20-40%.

9. The high strength dual-phase TRIP steel of claim 8, wherein the volume fraction of the fine-grained austenite contained in the retained austenite is 67%, the volume fraction of the large-grained austenite contained in the retained austenite is 33%.

10. The high strength dual-phase TRIP steel of claim 6, wherein after the high strength dual-phase TRIP steel is deformed, the volume fraction of austenite drops to 30-60%, the volume fraction of martensite increases to 30-60% and the volume fraction of the ferrite remains still.

11. The high strength dual-phase TRIP steel of claim 10, wherein the volume fraction of austenite drops to 52.7%, the volume fraction of martensite increases to 28.1% and the volume fraction of the ferrite remains 19.8%.

12. The high strength dual-phase TRIP steel of claim 1, wherein the high strength dual-phase TRIP steel includes vanadium carbide precipitations with a size of about 20-30 nm.

13. A method for making the high strength dual-phase TRIP steel, comprising: providing ingots, the ingots comprising 8-12 wt. % Mn, 0.4-0.6 wt. % C, 1-3 wt. % Al, 0.5-1 wt. % V, and a balance of Fe;hot rolling the ingots to produce a plurality of thick steel sheets with a thickness of 3-6 mm, then the steel sheets being treated by an air cooling process;warm rolling the steel sheets at a temperature of about 150-200° C. with a thicknesses reduction of about 15-44%;cold rolling the steel sheets at a room temperature with the thicknesses reduction of about 7%;annealing the steel sheets at the temperature of about 625-660° C. for 10-300 min to form the high strength dual-phase TRIP steel.

14. The method for making the high strength dual-phase TRIP steel of claim 13, wherein during the hot rolling process, a starting hot rolling temperature is 1100-1300° C., and a finishing hot rolling temperature is 800-1000° C., the thickness of each steel sheet is 3-6 mm.

15. The method for making the high strength dual-phase TRIP steel of claim 13, wherein the steel sheets are finally water quenched to the room temperature after the anneal process.

16. The method for making the high strength dual-phase TRIP steel of claim 13, wherein the high strength dual-phase TRIP steel comprises 8-12 wt. % Mn, 0.4-0.6 wt. % C, 1-3 wt. % Al, 0.5-1 wt. % V, and a balance of Fe.

17. The method for making the high strength dual-phase TRIP steel of claim 16, wherein the high strength dual-phase TRIP steel comprises 10 wt. % Mn, 0.47 wt. % C, 2 wt. % Al, 0.7 wt. % V, and a balance of Fe.

18. The method for making the high strength dual-phase TRIP steel of claim 13, wherein the high strength dual-phase TRIP steel consists of ferrite and retained austenite phases, the retained austenite phase consists of fine-grained austenite with a grain size that is smaller than 1 μm, and large-grained austenite with a size larger than 10 μm.

19. The method for making the high strength dual-phase TRIP steel of claim 18, wherein a volume fraction of austenite contained in the high strength dual-phase TRIP steel before a tensile test is 70-90%, a volume fraction of ferrite contained in the high strength dual-phase TRIP steel before the tensile test is 10-30%, a volume fraction of the fine-grained austenite contained in the retained austenite is 60-80%, a volume fraction of the large-grained austenite contained in the retained austenite is 20-40%.

20. The method for making the high strength dual-phase TRIP steel of claim 13, wherein the high strength dual-phase TRIP steel includes vanadium carbide precipitations with a size of 20-30 nm.