Method of forming high-plasticity metastable nanotwinned structures in austenitic steel alloys

Abstract

A method of forming a ductile, work hardenable, metastable nanotwinned austenitic steel alloy. Briefly, high purity feedstock is melted to obtain a cast steel alloy ingot, which is homogenized and cold rolled to reduce the thickness, and the homogenized and cold rolled steel alloy ingot is recrystallized to eliminate any possible casting defects, and the recrystallized steel alloy is cold rolled again to form a nanotwinned austenitic structure. The nanotwinned austenitic steel forms nanometer scale martensite platelets along the twin boundaries during subsequent deformation, resulting in an austenite / martensite nanolaminate structure in the steel alloy that exhibits good plasticity.

Description

[0001] Cross Reference to Related Applications

[0002] This application is related to U.S. Patent Application 18 / 153,374, filed on January 12, 2023. The disclosure of the U.S. Patent Application is incorporated herein by reference. TECHNICAL FIELD

[0003] The present invention relates to the field of steel alloys. More particularly, the present invention relates to a steel alloy with high strength and excellent ductility. BACKGROUND

[0004] Steels have excellent mechanical properties and economic benefits, and are the cornerstone of modern industry. Improving the strength level of steel materials can promote the progress of related technologies. For example, the continuous improvement of the strength of steel cables can increase the span of cable-stayed bridges by several times. The production of high-strength steel has always been a key issue in the competition of the steel industry. However, as the strength increases, it is often accompanied by a serious decrease in ductility. For example, when the yield strength of conventional high-strength steel exceeds 1 GPa, the uniform elongation is often limited to less than 10%. Poor ductility can lead to difficulties in the process of forming and shaping, limiting the application of traditional high-strength steel.

[0005] Strength and ductility represent the ability of a material to withstand external loads without plastic deformation and to absorb excessive loads, respectively, and are two key criteria for evaluating the reliability of materials in engineering and manufacturing. With the advancement of technology, how to improve the strength and ductility of metal materials has been a long-term pursuit of material design. Since the 1980s, a lot of effort has been invested in refining the grain size to the nanoscale to prepare nanocrystalline alloys to achieve the effect of strengthening alloys, and their strength has been successfully improved by one level. However, nanocrystalline alloys still have limited ductility problems, resulting in poor uniform elongation, which also means that it is still difficult to achieve the practical application of nanocrystalline alloys.

[0006] Nanocrystalline strengthening is considered as a strategy to optimize the strength-ductility synergy due to the unique interaction between dislocations and twin boundaries; however, it cannot overcome the intrinsic conflict between strength and ductility, i.e., the so-called strength-ductility inversion. The strength-ductility inversion is rooted in the deformation mechanism dominated by dislocations in metallic alloys, and introducing high density of crystal defects can impede the generation of dislocations to enhance the strength, but also consume the ability of dislocation accumulation, thus limiting the subsequent work hardening of the alloy. Twin boundaries are considered as an ideal structure to improve the ductility of nanoscale structured alloys due to their unique deformation mechanism; in other words, twin boundaries not only impede the occurrence of dislocations, but also allow dislocations to slide on the twin plane. Nevertheless, the ductility of nanotwinned metals is still much lower than that of coarse-grained metals due to the strength-ductility inversion, because the plastic deformation of nanotwinned alloys is also affected by the interaction between dislocations and boundaries, and when the twin thickness is optimized to the nanoscale, the ability of dislocation generation and accumulation is also limited.

[0007] To overcome the technical obstacles related to the strength-ductility inversion, it is necessary to provide an additional stress hardening capability in nanoscale structured alloys to enable the deformation mechanism of stress dislocations. Therefore, the present application solves this need. SUMMARY

[0008] One of the purposes of the present application is to provide a steel alloy that can overcome the strength-ductility inversion. More particularly, the steel alloy has a high plastic dynamically transformable nanotwinned (DT-NT) structure that not only prevents dislocations from strengthening the alloy, but also dynamically transforms into a nanolayered martensite / austenite dual-phase structure to increase the work hardening rate.

[0009] According to a first aspect of the present application, a method for forming ductile, work hardenable metastable nanotwins in an austenitic steel alloy, comprising the steps of:

[0010] melting high purity raw materials to obtain a cast steel alloy ingot;

[0011] homogenizing the cast steel alloy ingot to obtain a homogenized austenitic steel alloy ingot;

[0012] cold rolling the homogenized austenitic steel alloy ingot at room temperature to a reduction of 40-60%;

[0013] recrystallizing the cold-rolled austenitic steel alloy ingot to eliminate casting defects and form an as-recrystallized austenite (AR) steel alloy with a single face-centered cubic (fcc) structure of recrystallized grains; and

[0014] The AR steel alloy is cold-rolled to a reduction of 40-80% to form a metastable nanotwin (DT-NT) structure.

[0015] According to one embodiment of the present application, the metastable nanotwin structure provides twin boundaries as a platform to prevent dislocation generation and to partially transform austenite phase into martensite phase to form a nanolaminate dual phase structure of martensite / austenite upon stress to the alloy to exhibit good plasticity.

[0016] According to one embodiment of the present application, the method can further comprise a low temperature annealing step to promote dislocation recovery and stress relief.

[0017] According to one embodiment of the present application, the recrystallization annealing temperature is higher than an austenitizing temperature, wherein the austenitizing temperature can be between 1000-1200°C.

[0018] According to another embodiment of the present application, the method can further comprise a hot rolling or hot forging process at a temperature higher than the austenitizing temperature.

[0019] According to one embodiment of the present application, the high purity raw material comprises 22-26 wt% nickel, 0.8-2.5 wt% aluminum, 0.8-2.5 wt% silicon, 0.2-0.6 wt% carbon and 66.2-68.4 wt% iron.

[0020] According to another embodiment of the present application, the high purity raw material comprises 22-25 wt% nickel, 0.8-3 wt% silicon, 0.2-0.6 wt% carbon and 71.4-77 wt% iron.

[0021] According to a second aspect of the present application, a metastable nanotwin austenitic steel alloy with ductility, chromium-free and work hardenable, has a metastable nanotwin structure with a martensite phase extending from the nanotwin regions about 6-30 volume percent in the nanostructure, and the martensite phase forms a structure of martensite thin layers alternating with austenite thin layers; the nanotwin austenitic structure not only strengthens the alloy, but also promotes the martensite phase transformation to improve the work hardening ability, while improving its strength and ductility.

[0022] According to another embodiment of the present application, the metastable nanotwin austenitic steel alloy with ductility, ductility and work hardenable has a yield strength of at least about 1.4 GPa and an elongation of at least about 40%.

[0023] According to one embodiment of the present application, the steel alloy consists of 22-26 weight percent nickel, 0.8-2.5 weight percent aluminum, 0.8-2.5 weight percent silicon, 0.2-0.6 weight percent carbon, and 66.2-68.4 weight percent iron.

[0024] According to another embodiment of the present application, the steel alloy consists of 22-25 weight percent nickel, 0.8-3 weight percent silicon, 0.2-0.6 weight percent carbon, and 71.4-77 weight percent iron.

[0025] According to a third aspect of the present application, a safety part for an automobile prepared from the ductile work hardenable metastable nanotwinned austenitic steel alloy.

[0026] According to one embodiment of the present application, the safety part comprises a front longitudinal beam, a floor side reinforcement, a stationary interior part, a rear side beam, a B-pillar reinforcement, a roof bow, and an A-frame reinforcement.

[0027] According to a fourth aspect of the present application, a construction material prepared from the ductile work hardenable metastable nanotwinned austenitic steel alloy.

[0028] According to one embodiment of the present application, the construction material comprises a cable, a steel beam, and a scaffold.

[0029] According to a fifth aspect of the present application, an aircraft material prepared from the ductile work hardenable metastable nanotwinned austenitic steel alloy. BRIEF DESCRIPTION OF DRAWINGS

[0030] The following drawings describe embodiments of the present application in more detail, in which:

[0031] Figures 1A-1C shows the microstructure of a recrystallized austenite (AR) steel; Figure 1A is an electron backscattered diffraction (EBSD) map, Figure 1B is a phase distribution map, Figure 1C is an X-ray map showing a single bcc phase recrystallized microstructure;

[0032] Figures 2A-2D shows the mechanical properties of a DT-NT steel alloy; Figure 2A shows the engineering stress-strain curve of a DT-NT steel, and the stress-stress curve of a coarse grained AR steel for comparison; Figure 2B shows the tensile fractography of a DT-NT steel; Figure 2Cshows the relationship between the change in yield strength and the change in uniform elongation for the DT-NT steel of the present invention, and includes comparative data for various nanotwinned (NT) alloys, including (NT-Cu) nanotwinned copper alloys, NT steels, and gradient nanotwinned alloys; Figure 2D shows a comparison of the tensile properties of the DT-NT steel to high strength steels, including NT, martensitic, and transformation induced plasticity steels;

[0033] Figures 3A-3H shows the microstructure of the DT-NT steel prior to tensile testing; Figure 3A shows a backscattered electron diffraction pattern of a nanotwinned bundle embedded in a highly dislocated austenitic matrix; Figure 3B is a backscattered electron diffraction phase map;

[0034] Figure 3C is an XRD pattern showing the DI-NT steel; Figure 3D is a transmission electron microscope (TEM) bright field image of nanotwins, with an inset selected-area diffraction pattern (SADP) to confirm the twin structure; Figure 3E is a high resolution TEM image of nanotwins; Figure 3F is a statistical distribution plot showing the thickness of the T / M lamellar structure; Figure 3G is an elemental distribution map reconstructed from 3D-APT; Figure 3H is a one-dimensional distribution plot;

[0035] Figures 4A-4J shows the microscopic mechanisms of plastic deformation in the DT-NT steel; Figure 4A shows true stress-strain curves for the DT-NT and AR steels; Figure 4B shows strain hardening curves for the DT-NT and AR steels; Figure 4C is an EBSD orientation map of nanotwins in the DT-NT steel at 20% true strain; Figure 4D shows the phase distribution of nanotwins in the DT-NT steel alloy and in the AR alloy at 20% true strain; Figure 4E is an EBSD orientation map of nanotwins in the DT-NT and AR steels after fracture; Figure 4F shows the phase distribution of nanotwins in the DT-NT and AR steels after fracture; Figure 4G shows a TEM bright field image of nanotwins in the DT-NT steel at 20% true strain; Figure 4H is a selected-area diffraction pattern (SADP) showing the presence of martensite in nanotwins; Figures 4I-4JTEM dark field images showing twin / matrix lamellae and martensite, respectively, as the orientation between the martensite phase and the austenite phase follows the double Nishiyama-Wasserman orientation relationship;

[0036] Figures 5A-5F EBSD, TEM and XRD maps of the AR alloy to show the microstructural changes in the AR alloy during tensile deformation; Figure 5A is an EBSD map of the AR alloy before tensile deformation; Figure 5B is an EBSD map of the AR alloy after fracture;

[0037] Figure 5C is a TEM bright field image of the AR alloy at 5% strain; Figure 5D is a TEM bright field image of the AR alloy at 20% strain; Figure 5E is a TEM bright field image of the AR alloy after fracture; Figure 5F are XRD maps of the AR alloy before and after tensile deformation, showing that the volume fraction of martensite in the AR alloy after tensile deformation is about 7%;

[0038] Figures 6A-6C shows the deformation mechanism of the DT-NT steel at 5% strain; Figure 6A is an XRD map of the DT-NT steel alloy that still has a single face-centered cubic phase structure; Figure 6B is a representative TEM bright field image of the dislocation structure of the DT-NT steel; Figure 6C is a representative TEM bright field image of the nanotwin structure of the DT-NT steel;

[0039] Figure 7 shows the XRD patterns of the DT-NT steel at different engineering strains;

[0040] Figures 8A-8B shows the deformation mechanism of the dislocation region, where the inset is to show that no martensite phase transformation occurs in the dislocation region; Figure 8A is a TEM bright field image of the dislocation region at 20% strain; Figure 8B is a TEM bright field image of the dislocation region after fracture;

[0041] Figure 9 shows the nanohardness of the nanotwins (NT), nanolamellar dual-phase structure, and martensite; and

[0042] Figures 10A-10C shows the mechanism of the martensite phase transformation in the nanotwins; Figure 10A is a high-resolution TEM image showing the nucleation of the martensite phase at the twin boundary; Figure 10Bis a magnified TEM image of a twin boundary showing half-Shockley partial dislocation dipoles sliding on the twin plane, causing a martensitic phase transformation in the upper region; Figure 10C is a schematic illustration of the deformation mechanism of DT-NT steel. DETAILED DESCRIPTION

[0043] In the following description, ductile and work hardenable nanostructured steel alloys having metastable nanotwins are used as preferred examples. It will be apparent to those skilled in the art that modifications, including additions and / or substitutions, can be made without departing from the spirit and scope of the application. Specific details can be omitted in this description, so as not to obscure the teachings of the present disclosure; however, the description is made such that one skilled in the art could practice the teachings of the present disclosure without undue experimentation.

[0044] According to a first aspect of the present application, a method of forming a ductile, work hardenable metastable nanotwin structure in an austenitic steel alloy is provided. There are two basic principles for making a steel alloy with nanotwins, austenite / martensite nanolaminate nanostructure, including (1) tuning the stability of austenite to ensure that the steel alloy is mainly composed of austenite phase in the as-cast condition after water cooling; and (2) tuning the stacking fault energy of austenite during the forming process to introduce high density metastable nanotwins by deformation processing.

[0045] The method mainly comprises the following steps:

[0046] melting high purity raw materials to obtain a cast steel alloy ingot;

[0047] homogenizing the cast steel alloy ingot to obtain a homogenized austenitic steel alloy ingot;

[0048] cold rolling the homogenized austenitic steel alloy ingot at room temperature to a reduction of 40-60%;

[0049] recrystallizing the cold-rolled austenitic steel alloy ingot to eliminate casting defects and form an as-recrystallized austenite (AR) steel alloy having a single face-centered cubic (fcc) structure with recrystallized grains; and

[0050] cold rolling the AR steel alloy to a reduction of 40-80% to form a metastable nanotwin (DT-NT) structure

[0051] In an embodiment, the metastable nanotwin structure provides twin boundaries as a platform to prevent dislocation generation and partially transform the austenite phase into martensite phase when the alloy is subjected to stress, forming a nanolaminate dual-phase structure of martensite / austenite.

[0052] In one embodiment, the recrystallization annealing temperature is higher than the austenitization temperature, wherein the austenitization temperature can be in the range of 1000-1200 °C.

[0053] In one embodiment, the recrystallization annealing step can be replaced by a hot rolling or hot forging process, wherein the hot rolling or hot forging temperature is higher than the austenitization temperature.

[0054] In one embodiment, the high purity raw material comprises 22-26 wt% nickel, 0.8-2.5 wt% aluminum, 0.8-2.5 wt% silicon, 0.2-0.6 wt% carbon, and 66.2-68.4 wt% iron. To ensure the uniformity of the composition, all samples were repeatedly melted at least 5 times before casting into a rectangular copper mold with a thickness of 10 mm and a width of 12 mm. After homogenizing the ingot at 1050-1100 °C for 2 hours, the samples were cold rolled at room temperature to reduce the thickness by 60% (about 6 mm), and then recrystallized at 1050-1100 °C for 2 hours. Then, the samples were again cold rolled at room temperature to reduce the thickness by 75% (about 1.44 mm), and annealed at 170 °C for 12 hours. The prepared DT-NT steel samples exhibit an ultra-high strength of 1.4 GPa, which is almost five times that of the recrystallized steel (289 MPa), while maintaining a uniform elongation of 44%, showing that the prepared alloy is an ultra-high strength alloy with high ductility.

[0055] In another embodiment, the high purity raw material comprises 22-25 wt% nickel, 0.8-3 wt% silicon, 0.2-0.6 wt% carbon, and 71.4-77 wt% iron. To ensure the uniformity of the composition, all samples were repeatedly melted at least 5 times before casting into a rectangular copper mold with a thickness of 10 mm and a width of 12 mm. After homogenizing the ingot at 1050-1100 °C for 2 hours, the samples were cold rolled at room temperature to reduce the thickness by 60% (about 6 mm), and then recrystallized at 1050-1100 °C for 2 hours. Then, the samples were again cold rolled at room temperature to reduce the thickness by 75% (about 1.44 mm), and annealed at 170 °C for 12 hours. The prepared DT-NT steel samples exhibit a yield strength of 1320 MPa, which is more than five times that of the recrystallized steel (257 MPa), while maintaining a uniform elongation of 44%.

[0056] By combining cold rolling and low temperature annealing, the yield strength of the DT-NT steel can exceed 1.4 GPa, and the uniform elongation can be 40%. The DT-NT steel is superior to most commercially available high strength steels, and the excellent strength-ductility synergy of the DT-NT steel means that it can provide higher safety and has great potential in weight reduction and energy saving, while having higher formability.

[0057] Martensitic transformation, i.e. the partial transformation of austenite to martensite, is rarely observed in severely deformed alloys. Extensive studies have shown that a large amount of pre-deformation can increase the stability of austenite and suppress the martensitic transformation, i.e. the so-called mechanical stability of austenite, which is due to the high density of crystal defects (dislocations and grain boundaries) caused by pre-deformation, which destroys the orientation between martensite and parent austenite, inhibiting the nucleation of martensite; at the same time, the hardening of austenite will increase the strain energy required to overcome the transformation, thereby increasing the energy required for martensite nucleation. For example, the martensitic transformation start temperature (MS) of iron-9% manganese alloy will decrease with the increase of cold rolling reduction; in the present invention, the suppression of martensitic transformation occurs in the dislocation zone (non-nanotwinned zone). The mechanical stability of austenite is one of the main reasons for the strength-ductility inversion relationship in metastable alloys; however, in the DT-NT steel alloy of the present invention, the martensitic transformation in the nanotwins is significantly enhanced, providing a potential strategy for realizing the transformation induced plasticity (TRIP) effect in nanostructured alloys. As shown in Figure 4, the martensite originates from the nanotwins to form a layered structure of martensite / austenite, which produces high strength while maintaining high ductility.

[0058] According to a second aspect of the present invention, there is provided a high plastic metastable nanotwinned austenitic steel alloy; more particularly, the steel alloy can be prepared by the method described above and has a dynamically transformable nanotwinned austenitic structure, when the steel alloy is subjected to strain, part of the austenite in the structure can be transformed into martensite to form a nanolayered austenite / martensite dual-phase structure.

[0059] In one embodiment, the DT-NT steel alloy is composed of 22-26 weight percent nickel, 0.8-2.5 weight percent aluminum, 0.8-2.5 weight percent silicon, 0.2-0.6 weight percent carbon and 66.2-68.4 weight percent iron; in another embodiment, the DT-NT steel alloy is composed of 22-25 weight percent nickel, 0.8-3 weight percent silicon, 0.2-0.6 weight percent carbon and 71.4-77 weight percent iron.

[0060] To achieve low-cost industrial production, the chemical composition of the DT-NT steel is simplified to include only four to five elements, including nickel, aluminum, silicon, carbon, and iron. In the DT-NT steel, the addition of nickel, carbon, and silicon can effectively stabilize austenite and lower the martensitic transformation start temperature; however, excessive nickel, silicon, and carbon can make the austenite too stable and inhibit the deformation of nanotwin and the occurrence of martensite induced by deformation; the stacking fault energy can be adjusted by adding aluminum and carbon. Specifically, the addition of aluminum can increase the stacking fault energy, while the addition of carbon can reduce the stacking fault energy; in addition, the addition of aluminum can enhance the oxidation resistance and reduce the grain coarsening temperature, as well as reduce the density and increase the strength-to-weight ratio; in addition, the addition of silicon can increase the strength of austenite and enhance the oxidation resistance and corrosion resistance; however, excessive aluminum and silicon can promote the formation of brittle phases, reducing the ductility, toughness, and processability of the steel; on the other hand, the carbon element can effectively strengthen the austenite and martensite; however, excessive carbon can reduce the welding performance.

[0061] The alloy can be free of chromium, which reduces the cost of the steel alloy.

[0062] The nanotwinning austenite / martensite strengthening strategy can be applied to an austenitic steel (Fe-22% nickel-4% aluminum-4% silicon-2% carbon, atomic percentage). Nickel, silicon, and carbon atoms are austenite stabilizers that suppress the martensitic transformation during quenching, so after homogenization annealing and quenching, a single face-centered cubic (fcc) structure with recrystallized grains (average grain size of 70 μιη) can be obtained, and is considered a recrystallized (AR) steel alloy. Referring to Figures 1A-1C , respectively, showing the backscattered electron diffraction orientation map, phase distribution map, and X-ray map of the microstructure of the recrystallized (AR) steel alloy. By performing a controlled cold rolling process at room temperature, a nanotwin structure is introduced into the AR alloy while maintaining the fcc single phase, thereby preparing a DT-NT steel. In other words, by utilizing nanotwinning austenite and martensite strengthening, the effect of simultaneously improving strength and ductility is achieved.

[0063] Recrystallized DT-NT steel exhibits a similarly favorable strength-ductility synergy to commonly used commercial austenitic steels. Furthermore, after forming, the strength of DT-NT steel can be increased from 250–290 MPa to 1.3–1.4 GPa, a nearly fivefold increase, without sacrificing ductility, a property never before seen in commercially available steels. For many commercially available high-strength steels, the yield strength exceeds 1 GPa, resulting in a limited uniform elongation, sometimes less than 10%. Furthermore, compared to 304 and 316L stainless steels, DT-NT steel is safer, cheaper, and easier to process, making it highly competitive in the market. Furthermore, DT-NT steel can be prepared using conventional heat treatment and forming techniques, enabling low-cost, large-scale production. For these reasons, DT-NT steel is a potential candidate to replace the austenitic steels currently used extensively in automotive, construction, and aerospace applications.

[0064] In one embodiment, the DT-NT steel has an austenite / martensite structure exhibiting a nano-laminated martensite / austenite dual-phase microstructure, and the austenite-martensite lamellae are formed along the nano-twin boundaries, thereby forming an austenite / martensite nano-laminated structure in the steel alloy; in another embodiment, the amount of martensite phase present in the DT-NT steel is 7 to 30 volume percent.

[0065] In summary, the yield strength of DT-NT steel far exceeds that of other steels (such as 304 stainless steel, duplex steel, martensitic steel and high manganese steel), and maintains a considerable uniform elongation, which means that DT-NT steel is expected to extend the life of the final product and reduce its weight; at the same strength level, the uniform elongation of DT-NT steel is also one level higher than that of commercially available steel (such as martensitic steel), indicating that DT-NT steel is easier to form and can be prepared into complex shapes; in addition, since the production line of DT-NT steel is based on a combination of traditional technologies, DT-NT steel can be produced in an industrialized manner.

[0066] Example

[0067] Example 1

[0068] refer to Figures 2A-2D , the mechanical properties of DT-NT steel were further evaluated. Figure 2A As shown in Figure 2, by comparing the engineering stress-strain curves of recrystallized (AR) steel and DT-NT steel, it can be seen that the yield strength (σ_y) of DT-NT steel is 1439 MPa, which is almost five times that of AR steel (289 MPa), and DT-NT steel has excellent uniform ductility (44%), which is slightly better than that of CG sample (39.8%). Figure 2B As shown in Figure 2, the fracture surface of DT-NT steel has many dimples, indicating that it is a ductile fracture. Figure 2CAs shown, NT alloys, such as NT Cu, 316L and 304 stainless steels, and gradient nanotwinned (GNT) steels, exhibit a significant decrease in ductility accompanied by a multiplication of yield strength. Generally, when the yield strength of nanotwinned steels reaches gigapascal, the uniform elongation is completely lost; however, the superior strength-ductility synergy of DT-NT steels is superior to many high-strength steels, such as nanotwinned (NT) steels, martensitic steels, and transformation induced plasticity (TRIP) steels. As shown, at the same strength level, the uniform elongation of DT-NT steels is one order of magnitude higher than that of martensitic steels; meanwhile, the yield strength of DT-NT steels is also much higher than other steels with comparable uniform elongation. Therefore, DT-NT steels successfully overcome the inverse relationship between strength and ductility, which is contrary to the conventional understanding that cold rolling produces higher strength but severely sacrifices ductility. Figure 2D As shown, at the same strength level, the uniform elongation of DT-NT steels is one order of magnitude higher than that of martensitic steels; meanwhile, the yield strength of DT-NT steels is also much higher than other steels with comparable uniform elongation. Therefore, DT-NT steels successfully overcome the inverse relationship between strength and ductility, which is contrary to the conventional understanding that cold rolling produces higher strength but severely sacrifices ductility.

[0069] Example 2

[0070] To elucidate the principle of ultrahigh strength, microstructural studies of DT-NT steels were performed. A high-density nanotwinned fcc single-phase structure was found in the backscattered electron diffraction pattern of DT-NT steels (as shown in FIG. 2). Figures 3A-3B

[0071] Phase qualitative analysis was performed by using a Rigaku X-ray diffraction (XRD) instrument with Cu-Ka radiation and a monochromator. As shown in FIG. 3, the XRD pattern of DT-NT steels shows that no martensite phase appears during cold rolling due to its adiabatic heating effect. Figure 3C

[0072] In addition, the microstructural evolution during tensile testing was examined by transmission electron microscopy (TEM). TEM observations were performed on a 200 kV JEOL 2100F and a 300 kV double aberration-corrected TEM (Titan Cubed Themis G2300), and the samples to be examined were first mechanically polished to a thickness of 50 pm and then thinned by ion milling. As shown in FIG. 4, by TEM, it was found that DT-NT steels are composed of two types of substructures: twin bundles (indicated by dashed lines in the figure, hereinafter referred to as nanotwinned austenite) embedded in a deformed matrix of high-density dislocations (2.374 x 1014 m 15 m -2 ). Figure 3D Figure 3E By high-resolution TEM (HRTEM), it was found that the thickness of the twin / matrix lamella is in the nanometer range, from 3 nm to 150 nm, with an average value of about 22 ± 3.3 nm (as shown in FIG. 5); and by statistical measurement by TEM, it was found that the nanotwinned austenite accounts for 66% of the total volume. Figure 3F ​​​​

[0073] In addition, the elemental distribution was measured by 3D atom probe tomography (APT). APT measurements were performed in a voltage mode of 70k with a pulse repetition rate of 200kHz, a pulse fraction of 20%, and an evaporation detection rate of 0.2% atoms / pulse. The needle-shaped samples required for APT were prepared on a dual-beam focused ion beam (FIB) instrument (FEI Scios). Figures 3G-3H As shown in the figure, the chemical composition analysis results of DT-NT steel show that its elements are evenly distributed and no precipitation of carbides etc. is formed.

[0074] Example 3

[0075] Twin boundaries effectively block dislocations, resulting in high-strength properties.

[0076] According to Taylor's hardening law and the law of mixing, the yield stresses generated by nanotwin strengthening and dislocation strengthening are approximately 898 and 303 MPa, respectively. The strength of nanotwins reaches 1.36 GPa, which is still lower than the value calculated according to the Hall-Petch relationship (1.8 GPa).

[0077] Example 4

[0078] The uniqueness of DT-NT steel lies in that it not only strengthens the alloy like NT, but also improves its ductility by enhancing work hardening ability. Therefore, it is more important to deeply understand and study the micromechanism of plastic deformation of DT-NT steel. In short, in order to test its tensile properties, a wire-cut electric discharge machine was used to prepare a 12.5 mm gauge length and a cross section of 3.2 x 1.4 mm. 2 The dog-tooth shaped samples were used, and the loading direction of the tensile samples was aligned with the rolling direction. All tensile samples were ground with fine silicon carbide paper up to 2000#. The tensile tests were carried out in a Materials Testing System (MTS, Alliance RT30) tensile machine at room temperature with a strain rate of 2×10 -4 s -1 , Each sample was tested three times to ensure data reproducibility.

[0079] The uniqueness of nano-twinned austenite / martensite is that it not only strengthens the alloy like NT, but also improves its ductility by enhancing the work hardening ability. Figures 4A-4B As shown in the figure, unlike AR alloys which show a continuous decrease in work hardening energy, DT-NT steels show a distinct hardening stage, which allows the true tensile stress to be doubled to 2.1 GPa without sacrificing ductility. Generally speaking, a linear decrease in work hardening rate with increasing true strain is a typical characteristic of alloys, which is caused by the exhaustion of dislocation multiplication capacity.Figures 5A-5B )、TEM( Figures 5C-5E ) and XRD( Figure 5F ) shows the dominant deformation mechanism in AR alloy dislocations; e.g. Figures 5A-5B As shown in Figure 2, the orientation distribution diagrams of AR alloy before and after stretching are shown. The density of slip bands increases with the increase of strain, and only a small amount of martensite is observed in the AR alloy after fracture. Figures 5C-5E TEM bright field images show the deformation mechanism of AR alloy at 5% and 20% strain and after fracture, that is, the dislocation density increases with the increase of strain, and as shown in the inserted selected area electron diffraction (SAED) pattern, it is obvious that no martensitic transformation occurs in the dislocation structure; see also Figure 5F , in the AR sample after fracture, only a small amount of martensite (about 7 volume percent) was detected, which was not enough to improve the work hardening ability.

[0080] In contrast, DT-NT steel exhibits a multi-stage work hardening phenomenon. During the tensile process, pre-existing dislocations are first activated to maintain strain; dislocations detach from the pinning points on the solution atoms in the interstitial space, resulting in the emergence of upper and lower yield points. Figures 6A-6C As shown in XRD ( Figure 6A ) and TEM observation ( Figures 6B-6C ), when the strain is below 17% (for example, 5%), the deformation of DT-NT steel is mainly dominated by dislocation activity, and martensite rarely appears. Compared with AR alloys, the activation and dynamic recovery of pre-existing dislocations can cause a positive but lower work hardening rate; however, with further deformation (strain ≥ 17%), the work hardening rate of DT-NT steel increases significantly, and exceeds the work hardening rate of AR alloy at 20%, showing a clear work hardening stage; in addition, DT-NT steel can maintain a high work hardening rate (about 2.5GPa) until the true strain reaches 30%. This is completely different from the deformation softening in NT alloys, nanograined alloys or alloys after deformation treatment. It can be seen that the key reason for the excellent ductility of traditional NT alloys is the significant work hardening performance when the strain is greater than 17%.

[0081] In order to determine the reason for the high work hardening rate of DT-NT steel at strains greater than 17%, a systematic study of the microstructure at different strains was conducted. The results showed that when the true strain increased to more than 17%, martensite was produced; the volume percentage of martensite increased from 6 volume percent at a strain of 20% to 30 volume percent (according to Figure 7 Calculated from the XRD pattern in ). It can be seen that martensitic transformation is the main reason for the increase in work hardening rate; In addition, through EBSD ( Figures 4C-4F) and TEM Figures 4G-4J Further investigation of the microstructure revealed that the deformation-induced martensitic transformation mainly occurred in the nanotwins, and almost never in the dislocation regions Figures 8A-8B The statistical results from EBSD showed that the volume percentage of martensite in the nanotwins increased from 9.4 vol% at a true strain of 20% to 54.3 vol% at fracture, further confirming that almost all the martensitic transformation occurred in the nanotwin grains.

[0082] As shown in Figures 4C-4D , in the nanotwin regions, the martensite was lamellar and parallel to the twin lamellae. The selected area electron diffraction (SAED) Figure 4G ) of the nanotwin regions Figure 4H exhibited a double Nishiyama-Wasserman relationship between the martensite and austenite showed that the habit plane of the martensitic transformation was the twin plane. Under TEM dark field, the twins / matrix and the martensite could be distinguished, as shown in Figure 4I and Figure 4G . The nanoscale twin lamellae were shown to become irregular Figures 4I-4J ) due to the martensitic transformation, and with increasing strain, the lamellar martensite grew along the twin / matrix, transforming the austenitic nanotwins into a lamellar austenite / martensite dual-phase structure Figures 4E-4F .

[0083] By providing sufficient work hardening capacity, the introduction of high-plastic metastable nanotwins can serve as a way to provide ductility. Restricted by the parent nanotwins, the martensite also exhibited nanoscale thickness, thus forming a nanolamellar austenite / martensite dual-phase structure. As shown in Figure 9 , the newly generated nanodual-phase structure exhibited a nanohardness of 7.5 ± 0.3 GPa (the nanohardness of the martensite reached 8.1 ± 0.46 GPa), which was much higher than that of the parent nanotwins (5.6 ± 0.4 GPa), thus, at a larger strain, the continuous martensitic transformation in the nanotwins could dynamically strengthen the alloy and suppress strain localization, resulting in excellent ductility. This deformation mechanism of the DT-NT steel was quite different from that of the traditional NT alloy, in which the twin boundaries were stable during plastic deformation and no phase transformation occurred; however, with increasing strain, the ability of dislocation increase was soon exhausted, making shear bands and de-twinning dominate the local deformation, which seriously reduced the strengthening effect of the nanotwins, and thus led to early necking.

[0084] Example 5

[0085] Further relevant theoretical and experimental research is needed to explore the potential mechanism of martensitic transformation in nanotwins.

[0086] By high-resolution transmission electron microscopy (HRTEM), the twin boundaries provide nucleation sites for martensitic transformation ( Figure 10A ). It can be achieved through a two-step Shockley partial dislocation The movement transforms the face-centered cubic austenite into the body-centered tetragonal (bct)-α′ martensite. In the first step, every other (111) γ The atomic plane slips a Shockley partial dislocation to form a hexagonal close-packed (hcp)-ε phase; then, the half-Shockley partial dislocation dipole slides on each (0001)ε plane in the hcp phase, transforming the hcp-ε phase into bct-α′ martensite. Figure 10B As shown, with the assistance of twin boundaries, the martensitic transformation can be completed in only one step: the semi-Shockley partial dislocation The dipoles slide on the twin planes and are accompanied by atomic rearrangement to transform the twin planes into martensitic cores ( Figure 10C ), therefore, the twin boundary is the preferred site for martensite nucleation; in addition, the presence of many steps on the deformation-induced twin boundary promotes the nucleation of partial dislocations and accelerates the martensitic transformation.

[0087] In summary, nanotwinned austenite / martensite strengthening overcomes the inverted strength-ductility relationship of metallic materials, primarily by exploiting the dual functionality of twin boundaries: twin boundaries not only hinder dislocations to enhance strength, but also promote martensitic transformation, achieving the TRIP effect and enhancing ductility. This study demonstrates that by using nanotwins as preferential nucleation sites, martensitic transformation can be effectively promoted in severely deformed austenitic steels. This presents a new approach to optimizing the ductility of nanostructured alloys, and this simple and effective strengthening strategy is expected to be readily applicable to other metastable alloys.

[0088] The foregoing description has been presented for purposes of illustration and description of the present invention. It is not intended to be exhaustive or to limit the invention to the precise form disclosed, as many modifications and variations will be apparent to those skilled in the art.

[0089] The embodiments were chosen and described in order to best explain the principles of the invention and its practical applications, thereby enabling others skilled in the art to understand the invention for various embodiments and with various modifications as are suited to the particular use contemplated.

[0090] As used herein and unless otherwise defined, terms such as the terms "substantial," "substantially," "approximately," and "about," are used to describe and account for minor variations in, for example, measurements, values, amounts, sizes, etc. When used in conjunction with an event or circumstance, the terms can encompass instances that occur close to an exact occurrence of the event or circumstance. For example, when used in conjunction with a numerical value, these terms can encompass a range of variations less than or equal to ±10% of that numerical value, such as less than or equal to ±5%, less than or equal to ±4%, less than or equal to ±3%, less than or equal to ±2%, less than or equal to ±1%, less than or equal to ±0.5%, less than or equal to ±0.1%, or less than or equal to ±0.05%.

Claims

1. A ductile, chromium-free and work-hardenable metastable nanotwinned austenitic steel alloy, characterized in that, Comprising: a nanolayered martensite / austenite dual phase nanostructure, wherein the martensite phase extends from nanotwin regions and occupies 7 to 30 volume percent of the nanostructure, and the martensite phase forms a structure of alternating thin layers of martensite and austenite; wherein the thin layer thickness ranges from 3 nm to 150 nm, and the metastable nanotwinned austenitic steel alloy consists of 22-25 weight percent nickel, 0.8-3 weight percent silicon, 0.2-0.6 weight percent carbon, and 71.4-77 weight percent iron.

2. The ductile, chromium-free and work hardenable metastable nanotwinned austenitic steel alloy of claim 1, wherein the ductile, chromium-free and work hardenable metastable nanotwinned austenitic steel alloy has a yield strength of at least 1.4 GPa and an elongation of at least 40%.

3. The ductile, chromium-free and work hardenable iron-nickel-aluminum-silicon-carbon steel alloy of claim 1, wherein the, chromium-free and work hardenable metastable nanotwinned austenitic steel alloy further comprises aluminum and consists of 22-26 weight percent nickel, 0.8-2.5 weight percent aluminum, 0.8-2.5 weight percent silicon, 0.2-0.6 weight percent carbon, and 66.2-68.4 weight percent iron.

4. A safety part for a vehicle, characterized in that The safety components are made of the ductile, chromium-free and work hardenable metastable nanotwinned austenitic steel alloy of claim 1.

5. The security component of claim 4, wherein, The safety components include front longitudinal beams, floor side reinforcements, stationary interior components, rear side beams, B-pillar reinforcements, roof bows, and A-frame reinforcements.

6. A building material, characterized by The safety components are made of the ductile, chromium-free and work hardenable metastable nanotwinned austenitic steel alloy of claim 1.

7. The building material of claim 6, wherein, The construction materials include cables, steel beams, and scaffolding.

8. An aircraft material, characterized by The aircraft materials are made of the ductile, chromium-free and work hardenable metastable nanotwinned austenitic steel alloy of claim 1.

9. A method of making the ductile, chromium-free and work-hardenable metastable nanotwinned austenitic steel alloy of claim 1, characterized in that, Comprising: melting high purity raw materials to obtain an as-cast steel alloy ingot; homogenizing the as-cast steel alloy ingot to obtain a homogenized austenitic steel alloy ingot; cold rolling the homogenized austenitic steel alloy ingot at room temperature to a reduction of 40-60%; recrystallizing the cold rolled austenitic steel alloy ingot to eliminate casting defects and form a recrystallized austenitic steel alloy with a single face-centered cubic structure of recrystallized grains; and cold rolling the recrystallized austenitic steel alloy to a reduction of 40-80% to form a metastable nanotwinned structure, wherein the metastable nanotwins can gradually transform into an austenite / martensite nanolaminate structure during tensile deformation to exhibit good plasticity; wherein the martensite phase is present in an amount of 7 to 30 volume percent.

10. The method of claim 9, further comprising a low temperature annealing step to promote dislocation recovery and stress relief.

11. The method of claim 10, wherein the low temperature annealing is at a temperature between 140-200 °C.

12. The method of claim 11, further comprising a hot rolling or hot swaging treatment at a temperature above the austenitizing temperature.

13. The method of claim 9, wherein the high purity feedstock comprises 22 to 26 weight percent nickel, 0.8 to 2.5 weight percent aluminum, 0.8 to 2.5 weight percent silicon, 0.2 to 0.6 weight percent carbon, and 66.2 to 68.4 weight percent iron.

14. The method of claim 9, wherein the high purity feedstock comprises 22 to 25 weight percent nickel, 0.8 to 3 weight percent silicon, 0.2 to 0.6 weight percent carbon, and 71.4 to 77 weight percent iron.

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