Carbon-doped alloy steel with room-temperature super-ductility and preparation process thereof

Abstract

The invention relates to carbon-doped alloy steel with room-temperature super ductility. The carbon-doped alloy steel comprises the following chemical components in percentage by weight: 10 to 30 percent of nickel (Ni), 0.01 to 5 percent of silicon (Si), 0.01 to 2 percent of carbon (C) and 60 to 90 percent of iron (Fe). Different from a traditional super-ductility alloy which can only obtain superplasticity at a high temperature, a super-ductility steel is introduced, and the uniform elongation at the room temperature exceeds 100%. Meanwhile, different from the deformation softening of the super-ductility alloy, the super-ductility steel shows excellent work hardening capability, so that the super-ductility steel can obtain enough strength to be applied to various fields.

Description

CROSS-REFERENCE TO RELATED APPLICATIONS

[0001] This application claims priority to U.S. patent application No. 18 / 353,102 filed on July 16, 2023, the disclosure of which is incorporated herein by reference. Technical Field

[0002] The present invention generally relates to the field of materials science and technology. In particular, the present invention relates to the development of a new alloy steel with excellent ductility at room temperature. Potential applications of ultra-high ductility alloy steel include, but are not limited to, critical safety components such as front side support bars, floor side reinforcements, static interior components, rear side support bars, B-pillar reinforcements, roof crossbars, A-pillar reinforcements, etc. Background Art

[0003] In order to manufacture complex parts with high precision in a single machining process, it is crucial to use extremely ductile materials, such as super-ductile (defined as an extremely high ductility of more than 300%). 1 However, it is usually necessary to heat the system at a higher temperature (T / T m , T is the actual temperature, T m This ductility is achieved at temperatures below the melting point of the material and often exhibits limited uniform extension and deformation softening.

[0004] Traditionally, achieving this high ductility relies on triggering thermally activated deformation mechanisms such as grain boundary sliding and cobalt creep at high homogenous temperatures. 2，3 Conventional high-ductility alloys are not suitable for large-scale production due to their excessive energy consumption. In addition, the thermally activated deformation mechanism leads to deformation softening during superductility and insufficient strength for structural applications. 1，4 Therefore, alloys with ultra-high ductility and good strength at low temperatures (especially room temperature) must be developed to reduce carbon footprint and energy consumption.

[0005] The uniform elongation of metallic materials at low homologous temperatures depends on their ability to undergo sufficient strain hardening, which enables them to follow increasing flow stresses. 5 However, the strain hardening capacity decreases with increasing strain, mainly due to the gradual accumulation of dislocations, setting a fundamental limit to the room temperature ductility of metallic materials. 6 The main methods to improve room temperature ductility are to reduce the stacking fault energy and promote twinning induced plasticity (TWIP) or transformation induced plasticity (TRIP). 7 However, despite the presence of TWIP and TRIP, the room temperature uniform elongation of metal alloys rarely exceeds 100% 8 . Summary of the invention

[0006] Therefore, the object of the present invention is to provide a carbon-doped alloy steel having room temperature super ductility. In a first aspect, the carbon-doped alloy steel comprises the following chemical components: 10-30% by weight of nickel (Ni), 0.01-5% by weight of silicon (Si), 0.01-2% by weight of carbon (C) and 60-90% by weight of iron (Fe). The addition of carbon to the FeNiSi alloy results in a phase composition transformation from a single martensite phase to a single austenite phase.

[0007] According to one aspect of the present invention, optionally, the carbon-doped alloy steel further comprises 0.1-20% by weight of manganese (Mn).

[0008] According to one aspect of the present invention, the carbon-doped alloy steel has a uniform elongation of at least 100%.

[0009] According to one aspect of the present invention, the yield strength of the carbon-doped alloy steel after cold rolling is at least 1200 MPa.

[0010] According to one aspect of the present invention, the carbon-doped alloy steel elongates uniformly during the entire tensile test without obvious necking.

[0011] According to one aspect of the present invention, the carbon-doped alloy steel exhibits a face-centered cubic (fcc) structure with a fully recrystallized morphology.

[0012] According to one aspect of the present invention, the tensile strength of the carbon-doped alloy steel is increased by at least 30% relative to the yield strength.

[0013] According to one aspect of the present invention, the carbon-doped alloy steel exhibits a true stress in the range of 1500-2200 MPa when the true strain ranges from 80% to 100%.

[0014] The present invention is the first to observe the room temperature super ductility of FeNiSiC alloys. This excellent ductility is achieved through sustainable work hardening ability. Specifically, carbon plays a vital role in stabilizing the austenite phase, and FeNiSiC alloys undergo dynamic redistribution of carbon during deformation. This element redistribution leads to local changes in stacking fault energy and austenite stability, triggering the sequential activation of multiple deformation mechanisms in a single fcc phase FeNiSiC alloy, namely dislocation slip, deformation twinning, and martensitic phase transformation. These internal defects gradually refine the structure of FcNiSiC alloys from tens of microns to nanometers, thereby achieving sustainable work hardening ability. These findings provide theoretical and experimental basis for the strategy of integrating various work hardening mechanisms using the dynamic redistribution of interstitial solute elements to improve the ductility of metallic materials.

[0015] In a second aspect, the present invention provides a method for preparing a carbon-doped alloy steel having room temperature super ductility, comprising: arc melting a raw material in an argon environment after titanium adsorption to obtain an ingot; homogenizing the ingot; cold rolling the ingot at room temperature to reduce the thickness by 40-60%; recrystallizing the cold-rolled ingot to obtain a carbon-doped alloy steel; cold rolling the carbon-doped alloy steel to reduce the thickness by 50-80%; and annealing the carbon-doped alloy steel to adjust the grain size.

[0016] According to one embodiment, the homogenous reaction temperature is in the range of 1000-1150° C., and the reaction time is in the range of 1-2 hours.

[0017] According to one embodiment, the recrystallization reaction temperature is in the range of 1000-1150° C., and the reaction time is in the range of 1-2 hours.

[0018] According to one embodiment, the annealing reaction temperature is in the range of 800-1200° C., and the reaction time is in the range of 0.5-30 minutes.

[0019] The present invention has the following advantages: (1) Super ductility alloy steels exhibit a room temperature uniform elongation of over 100%. Compared with traditional super ductility that can only be achieved at high temperatures, this room temperature super ductility can effectively reduce the energy cost of the forming process. (2) Unlike the deformation softening observed during superductile deformation, superductile alloy steels exhibit excellent work hardening capabilities, enabling them to obtain sufficient strength after the forming process. Therefore, this superductile alloy steel is an excellent candidate for near-net forming. (3) Compared with many existing steels (such as iron-manganese steel, 304 steel and 316L steel), super-ductile alloy steels exhibit excellent strength-ductility synergy. BRIEF DESCRIPTION OF THE DRAWINGS

[0020] The present invention will be more readily understood from the following description of exemplary embodiments with reference to the accompanying drawings, in which:

[0021] Figure 1A A comparison of the engineering stress-strain curves of FeNiSi and FeNiSiC is depicted.

[0022] Figure 1B The true stress-strain curves of FeNiSi and FeNiSiC alloys are depicted.

[0023] Figure 1C Images of a FeNiSiC alloy tensile specimen before fracture (left) and after fracture (right) are shown.

[0024] Figure 2Tensile fractograms of FeNiSi and FeNiSiC alloys are depicted.

[0025] Figure 3A Electron backscatter diffraction (EBSD) inverse pole figure (IPF) diagram and phase distribution map are shown, showing the fully martensitic structure of the FeNiSi alloy.

[0026] Figure 3B Electron backscatter diffraction (EBSD) inverse pole figure (IPF) diagram and phase diagram are shown, showing the fully austenitic structure of the FeNiSiC alloy.

[0027] Figure 3C X-ray diffraction (XRD) patterns of FeNiSi alloy and FeNiSiC alloy are depicted.

[0028] Figure 3D Transmission electron microscopy (TEM) bright-field images of FeNiSiC alloy are shown.

[0029] Figure 3E The EDS image of the FeNiSiC alloy and the three-dimensional atom probe tomography (3D-APT) of FeNiSiC are shown, demonstrating the uniform distribution of the elements.

[0030] Figure 3F The corresponding one-dimensional (1-D) composition profiles quantitatively showing the element distribution along the length of the tip are depicted.

[0031] Figure 4A TEM image showing the microstructure of the FeNiSi alloy.

[0032] Figure 4B An energy dispersive spectroscopy (EDS) graph is shown, which illustrates the uniform distribution of each element in the carbon-free alloy.

[0033] Figure 5A The strain hardening curve of FeNiSiC alloy is depicted, showing three strain hardening stages.

[0034] Figure 5B The dynamic changes of the work hardening index (i.e., instantaneous work hardening index) of FeNiSiC alloy during the uniform deformation stage are depicted.

[0035] Figure 5C Depicted are XRD patterns showing the phase composition of the FeNiSiC alloy at different true strains.

[0036] Figure 5D The microstructure evolution of the FeNiSiC alloy with increasing strain is shown.

[0037] Figure 6Bright-field TEM image of a FeNiSiC alloy at 15% true strain showing dislocation tangles and dislocation cells.

[0038] Figure 7 shows bright-field TEM images of deformation twins in FeNiSiC alloy under 25% true strain.

[0039] Fig. 8A High-resolution TEM images of deformation twins in FeNiSiC alloy at 45% true strain are shown.

[0040] Figure 8B Shown is the fast Fourier transform of deformation twins in a FeNiSiC alloy at a true strain of 45%.

[0041] Fig.9A Bright-field TEM image of the martensite core in a FeNiSiC alloy at a true strain of 72% is shown.

[0042] Fig. 9B Shown is the Kurdjumov-Sachs orientation relationship in FeNiSiC alloy at 72% true strain. Dark-field TEM image of the martensite phase.

[0043] Fig. 9C Shows Fig.9A The corresponding selected area diffraction pattern in .

[0044] Fig.9D Shown is the Nishiyama-Wasserman orientation relationship in FeNiSiC alloy at 72% true strain Dark-field TEM image of the martensite phase.

[0045] Fig. 10A High-resolution TEM image of the martensite core in a FeNiSiC alloy at 72% true strain is shown.

[0046] Fig. 10B Shown is the FFT of the martensite core in a FeNiSiC alloy at 72% true strain.

[0047] Fig.11A A bright field TEM image of a carbon-doped alloy after fracture is shown. The inset shows the electron beam and

[011] fcc and

[001] bcc Tightly parallel selected area diffraction patterns.

[0048] Fig. 11B Shows Fig.11ADark-field TEM image of the austenite (fcc) phase.

[0049] Fig. 11C Shows Fig.11A Dark-field TEM image of the medium martensite (bcc) phase.

[0050] Fig.12 Shown is the dynamic carbon distribution during the room temperature tensile test.

[0051] Fig.13A Schematic diagram depicting the change in characteristic size as a function of true strain. Multiple strain hardening mechanisms lead to shortening of the mean free path of dislocations; and

[0052] Fig. 13B The tensile properties of existing carbon-doped alloy steels and other existing high-ductility metallic materials at room temperature are depicted. DETAILED DESCRIPTION

[0053] Steel accounts for about 65% of the average car's weight. In 2020, China's automotive industry's consumption of high-strength steel reached 52.5 million tons. Due to the limited ductility of existing steels, the design of mechanical parts often needs to be as simple as possible, even at the expense of practicality. In addition, complex mechanical parts often need to be manufactured separately and subsequently assembled, resulting in high energy and time costs.

[0054] Adding Ni, C, and Si to steel can effectively stabilize austenite and reduce the starting temperature of martensite transformation. However, excessive Mn, Ni, Si, and C make austenite too stable and inhibit the formation of deformation nanotwins and strain-induced martensite. Adding Si increases the strength of austenite and enhances its oxidation resistance and corrosion resistance. However, excessive Si promotes the formation of brittle phases and reduces the ductility, toughness, and machinability of steel.

[0055] The present invention has found that carbon (C) plays an important role in controlling the deformation mechanism by dynamically affecting the stacking fault energy and austenite stability. Super ductile carbon-doped steel makes near-net forming of parts of complex shapes possible. In particular, the present invention provides a carbon-doped alloy steel with room temperature super ductility. The carbon-doped alloy steel includes the following chemical components: 10-30% nickel (Ni) by weight, 0.01-5% silicon (Si) by weight, 0.01-2% carbon (C) by weight, and 60-90% iron (Fe) by weight. Due to the incorporation of C into the FeNiSi alloy, the carbon-doped alloy steel exhibits a room temperature uniform elongation of at least 100%.

[0056] The basic principles for designing superductile steels include two aspects: (1) controlling the stability of the austenite phase to ensure that the steel is mainly composed of austenite after quenching, and (2) adjusting the stacking fault energy and stability of austenite to promote the sequential activation of various deformation mechanisms such as dislocations, strain-induced twinning, and martensitic transformation.

[0057] In one embodiment, the carbon-doped alloy steel includes four elements: 10-30% nickel (Ni), 0.1-5% silicon (Si) by weight, 0.01-2% carbon (C) by weight, and 60-90% iron (Fe) by weight. In this article, "wt%" refers to weight percentage, which is an indicator of the relative weight of a specific element or component in the alloy steel relative to the total weight of the alloy steel itself.

[0058] Preferably, the carbon-doped alloy steel comprises 23-25% by weight of Ni, 0.8-3% by weight of Si, 0.1-0.5% by weight of C and 71.5-76.1% by weight of Fe.

[0059] In other embodiments, the carbon-doped alloy steel contains five elements: 10-30 weight % nickel (Ni), 0.1-20 weight % manganese (Mn), 0.01-5 weight % silicon (Si), 0.01-2 weight % carbon (C) and 43-89.88 weight % iron (Fe).

[0060] Preferably, the carbon-doped alloy steel includes 18-26 weight % Ni, 1-10 weight % Mn, 0.05-4 weight % Si, 0.05-9 weight % C and 59.1-80 weight % Fe.

[0061] In one embodiment, the carbon-doped alloy steel exhibits a uniform elongation of at least 100%. For example, the carbon-doped alloy steel exhibits a uniform elongation of at least 110%, at least 120%, at least 130%, at least 140%, at least 150%, at least 160% or at least 170%.

[0062] In one embodiment, the carbon-doped alloy steel has a yield strength of at least 1200 MPa after cold rolling. For example, the yield strength after cold rolling is at least 1300 MPa, at least 1350 MPa, at least 1400 MPa, which is three times the yield strength before cold rolling.

[0063] Preferably, with a thickness reduction of 76% after cold rolling, the carbon-doped alloy steel can achieve a good combination of strength and ductility, particularly a yield strength of up to 1400 MPa and a uniform elongation of 40%, which exceeds most existing Fe-Mn steels, 304 and 316L stainless steels.

[0064] In one embodiment, the tensile strength of the carbon-doped alloy steel is at least 30% higher than the yield strength. For example, the tensile strength is increased by at least 40%.

[0065] These exceptional forming capabilities and unique mechanical properties allow super ductile steels to be used in near-net-shape processes, where they can be formed into complex parts while achieving sufficient strength for practical applications.

[0066] In other aspects, the present invention also provides a method for preparing a carbon-doped alloy steel with room temperature super ductility, the processing of which includes four steps: (1) The alloy is prepared by arc melting a high purity (99.9 wt%) raw material in an argon environment with titanium adsorption. (2) In order to eliminate possible casting defects, the as-cast ingot is homogenized at 1000-1150°C for 2 hours, and then the ingot is cold rolled at room temperature to reduce the thickness by 40-60%, and then recrystallized at 1000-1150°C for 1-2 hours. This step can also be replaced by hot forging and hot rolling at a temperature above the austenitizing temperature. (3) The grain size is adjusted by cold rolling and subsequent recrystallization annealing. The sample is cold rolled to reduce the thickness by 50-80%, and then annealed at 800-1200°C for 0.5-30 minutes. The annealing parameters should be selected according to the thickness of the cold rolled sample.

[0067] Considering the excellent combination of strength and ductility exhibited by super-ductile alloy steels after forming, they have the potential to replace many existing advanced high-strength steels in applications such as automobiles, construction and shipbuilding. Super-ductile alloy steels are expected to attract great attention from automobile manufacturers and steel companies.

[0068] Example

[0069] Example 1

[0070] Preparation of super ductile alloys

[0071] Carbon is added to Fe-Ni-Si alloys with poor ductility to achieve room temperature superplasticity. In one embodiment, the phase composition of the super ductile alloy includes 22-24% Ni by weight, 1-2.5% Si by weight, 0.3-0.5% C by weight, and 73-76.7% Fe by weight. The material is then homogenized at 1050-1100°C for 2 hours, then cold rolled at room temperature to reduce the thickness by 40-60%, and then recrystallized at 1050-1100°C for 2 hours. Cold rolled at room temperature to reduce the thickness by 76%, and then annealed at 1100°C for 1 minute. The prepared super ductile alloy steel (FeNiSiC alloy) can exhibit an ultra-high uniform elongation of about 174%.

[0072] Example 2

[0073] Comparison between FeNiSi alloy and FeNiSiC alloy

[0074] In this example, a FeNiSiC alloy was prepared and processed according to the description in Example 1. Figure 1A Comparison of the engineering strain-stress curves of Fe-24Ni-4Si (referred to as FeNiSi) and Fe-24Ni-4Si-2C (referred to as FeNiSiC) shows that the two alloys exhibit significantly different mechanical properties. The FeNiSi alloy exhibits a limited uniform elongation of 3±0.1% and a total elongation of 10±1.1%. However, after carbon doping, the FeNiSiC alloy becomes significantly softer and exhibits superductility, with a uniform elongation of 174±4.1%, which is more than 50 times that of the FeNiSi alloy. Unlike high homologous temperature superplasticity characterized by work softening, the FeNiSiC alloy exhibits work hardening behavior. The tensile strength of the FeNiSiC alloy increases from 294±15.6MPa at the yield point to 735±14.4MPa before fracture. In addition, the FeNiSiC alloy exhibits higher true stress compared to the FeNiSi alloy. At a true strain of 103%, the true stress of the FeNiSiC alloy reaches up to 2000MPa, which is twice that of the FeNiSi alloy (1000MPa). Figure 1B shown).

[0075] Appearance of tensile specimens ( Figure 1C ) and fracture morphology ( Figure 2 ) showed that FeNiSiC elongated uniformly throughout the tensile test without significant necking. This room temperature superductility accompanied by work hardening behavior has not been reported in previous studies of steels or other superplastic alloys.

[0076] Example 3

[0077] Characterization of super ductile FeNiSiC alloy

[0078] The significant difference in tensile properties between FeNiSi and FeNiSiC alloys raises the question of the role of carbon in modifying the initial microstructure, thereby transforming the brittle alloy into a quasi-superplastic alloy. To address this question, the pre-deformed microstructure was systematically investigated using techniques such as electron backscatter diffraction (EBSD), X-ray diffraction (XRD), atom probe tomography (APT) and transmission electron microscopy (TEM).

[0079] The inverse pole figure (IPF) is a commonly used graphical tool in materials science and metallurgy to visualize the spatial distribution of crystal orientation and grain arrangement. Referring to Figure 3, the microstructures of FeNiSi and FeNiSiC alloys are shown. The FeNiSi alloy exhibits a single-centered cubic (bcc) phase ( Figure 3A and 3C ), and have equiaxed grains with a lamellar structure, e.g. Figure 3A In addition, TEM observation results confirm that the microstructure of FeNiSi alloy is composed of lath-shaped martensite, as shown in the IPF diagram. Figure 4A As shown. Bright field TEM images show lath-like martensite with high density of dislocations. In contrast, the FeNiSiC alloy shows a single fcc structure with a fully recrystallized morphology ( Figure 3B-3C ). Adding carbon to the FeNiSi alloy effectively transforms the phase composition from a single martensite phase to a single austenite phase. TEM bright field images of FeNiSiC alloys are shown in Figure 2. Figure 3D shown.

[0080] The element distribution of the two alloys was characterized by energy dispersive spectroscopy (EDS) on TEM. The results showed that all elements were uniformly distributed in the FeNiSiC alloy ( Figure 3E ) and FeNiSi alloys ( Figure 4B )middle.

[0081] Considering the limited accuracy of TEM-EDS in characterizing light elements such as carbon, the chemical composition of the FeNiSiC alloy was further investigated using atom probe tomography (APT), which confirmed the chemical homogeneity of the alloy with no carbon segregation inside the grains. Figure 3F The one-dimensional carbon composition profile along the needle tip axis shows that there is no significant fluctuation in carbon concentration. Carbon, as an important austenite stabilizer, effectively enhances the stability of austenite and reduces the martensitic transformation starting temperature (Ms), thereby inhibiting the martensitic transformation during heat treatment. Figure 5C XRD patterns are shown, showing the phase composition of FeNiSiC alloy at different true strains (5%, 15%, 25%, 45%, 72% and fracture). Carbon doping transforms the phase composition of FeNiSi alloy from brittle body-centered cubic (bcc) structure to ductile face-centered cubic (fcc) structure of FeNiSiC alloy. This transformation is the main reason for the super ductility of FeNiSiC alloy.

[0082] Unlike high-temperature superductility with high strain rate sensitivity, the huge uniform ductility of FeNiSiC alloys comes from the ability to sustain work hardening. Figure 5A , in the initial stage (stage I, at the true strain ε T Less than 15%), the work hardening rate drops sharply, and the reduction in the true strain εT Between 15% and 45% (Phase II) is inhibited. In addition, T ≥45% (stage III), the FeNiSiC alloy shows a continuous increase in work hardening ability until fracture (ε T =102%). Correspondingly, the work hardening exponent increases with the true strain and reaches an extremely high value of 1.2 when the true strain reaches 102%, showing the excellent work hardening ability of FeNiSiC alloy ( Figure 5B ).

[0083] Systematic microstructural characterization indicates that the multiple work-hardening behaviors in FeNiSiC alloy originate from the sequential activation of different deformation mechanisms, namely, dislocation activation, deformation twinning, and γ→α martensite. Similar to most polycrystalline FCC alloys, dislocations are the main carriers of plastic strain in the first stage.

[0084] Figure 5D and Figure 6 They show the T Typical dislocation network at 5% and at ε T The dislocation entanglement at 15%. As the strain increases, the dislocation density also increases, resulting in a decrease in the dislocation accumulation rate. The decrease in the dislocation accumulation rate is the reason for the decrease in the work hardening rate.

[0085] In stage II, further stretching leads to the activation of deformation twins in the FeNiSiC alloy. Figure 7 Shown is a transmission electron microscopy bright-field image of deformation twins in FeNiSiC alloy when the true strain reaches 25%. Fig. 8A shows a step at the deformation twin boundary at a true strain of 45%, and Figure 8B The corresponding Fast Fourier Transform (FFT) is shown. Fig.9A shows a high-resolution TEM image of the martensite core in a FeNiSiC alloy at 72% true strain, and Fig. 9B The corresponding FFT is shown.

[0086] For example, in ε T =45%, a high density of nanoscale deformation twins (average twins / matrix lamellae of 26 ± 6.4 nm) with stacking faults was observed, with a volume fraction of 30 ± 5% (estimated by TEM). The rotation observed in the selected area electron diffraction pattern indicates that the twin boundaries are distorted due to the strong interaction with the dislocations. The presence of twin boundaries hinders further slip of dislocations, thereby enhancing the work hardening ability of the material. This allows the material to maintain greater plasticity without necking.

[0087] After deformation twinning, α′-martensite transformation is also induced in the third stage to maintain plastic strain. The results show that α′-martensite tends to nucleate in the twin region and transform at the twin boundary according to the Kurdjumov-Sachs orientation relationship ( Fig. 9C ) nucleation, or in the twin / matrix lamellae according to the Nishiyama-Wasserman orientation relationship ( Fig.9D ) nucleation.

[0088] In addition, the volume fraction of martensite increases with increasing strain, from 20% by volume at a true strain of 72% to 34% by volume at failure. At the same time, the volume fraction of deformation twins increases to 80% by volume at failure. The large amount of deformation twins and strain-induced α′-martensite phase transformation in the third stage produce a nanoscale martensite-austenite dual-phase structure (Figures 10-11), resulting in a continuous increase in the work hardening rate in stage III.

[0089] The sequential activation of deformation twinning and α′-martensitic transformation is anomalous in single-phase fcc steels with homogeneous recrystallized structures. In particular, FeNiSiC alloys with stacking fault energies of 22–30 mJ are considered TWIP alloys. Typically, in metastable alloys, a layered structure is required to unify these two enhanced hardening mechanisms, where grain size stratification induces extensive variations in phase stability, promoting the simultaneous occurrence of different deformation mechanisms.

[0090] Chemical distribution plays a crucial role in the stability of austenite. In order to study the microstructural origin of various deformation mechanisms, the present invention uses APT-TEM combined technology to analyze and characterize the evolution of chemical distribution in FeNiSiC alloy during tensile testing.

[0091] Fig.12 The true strain is shown to be 45% ( Fig.12 (a)) and after fracture ( Fig.12 (b)) APT and TEM characterization of carbon distribution in FeNiSiC alloy. (a1) and (b1) are the bright field image, dark field image and selected area electron diffraction pattern of the APT probe, respectively. (a2) and (b2) are the corresponding carbon atomic distribution maps reconstructed using 3D-APT. (a3) ​​and (b3) are carbon isovalue curves with concentrations of 3.5% and 4% atomic percentage, respectively. (a4) and (b4) are the one-dimensional composition distributions of carbon along the directions indicated by the arrows in (a3) ​​and (b3), respectively.

[0092] Unexpectedly, the results showed that carbon, an important austenite stabilizing element, tends to dynamically redistribute during tensile testing at room temperature ( Fig.12 ). Analyze ε through APT T =45% containing nanotwins (a1). The C atomic map (a2) and the 3.5 atomic % carbon isoconcentration surface (a3) ​​reveal nanoscale carbon-rich clusters in the twinned region. The one-dimensional (1D) concentration distribution (a4) quantitatively reveals carbon fluctuations. The carbon-rich region contains 3 atomic % carbon, accompanied by carbon-poor regions with carbon concentrations as low as 1 atomic %. The redistribution of carbon becomes more obvious with increasing strain. In the FeNiSiC alloys after fracture (b1-b4), the carbon content of the retained austenite is as high as 5%, while the martensite shows a low carbon content of 1%.

[0093] TWIP and α′-martensitic transformation-induced plasticity have been extensively studied in many alloys. However, this has only been able to push the room temperature uniform elongation to a maximum of 103% and 50% ( Fig.13A ). Although the TWIP and α′-TRIP effects can be activated together in some heterostructured alloys, the total elongation is difficult to exceed 100%. The total elongation of the FeNiSiC alloy of the present invention completely exceeds that of various TWIP / TRIP Fe-Ni steels, Fe-Mn steels and fcc phase high entropy alloys (HEA). Specifically, the total elongation of the FeNiSiC alloy is 1.8 times that of Fe-20Mn-1.2C steel.

[0094] The obvious TWIP and TRIP effects in FeNiSiC alloy can be attributed to the following factors. First, the density of deformation twins and martensite is extremely high, dividing the coarse grains into nanostructures, namely, nanotwins with an average thickness of 25 ± 6.4 μm and nanoscale martensite / austenite dual-phase structure with an average martensite size of 13 ± 0.9 μm ( Fig. 13B and Figure 10). On the other hand, the dynamic distribution of carbon from fresh martensite to austenite increases the stability of austenite and reduces the transformability of retained austenite. The critical stress for martensitic transformation increases dynamically with increasing strain, allowing the martensitic transformation to proceed in a gradual and controllable manner. The martensitic transformation rate per engineering strain (%) is only 0.29% by volume, an order of magnitude lower than that of most TRIP steels, which helps to maintain sustainable enhanced hardening capacity under extremely high strain conditions (about 117% engineering strain).

[0095] In summary, the super ductile steel of the present invention is an excellent candidate for near net forming processes. For example, when cold rolled with a thickness reduction of about 76%, the super ductile steel exhibits a yield strength of about 1400 MPa and a uniform elongation of about 40%. These values ​​exceed those of most Fe-Mn, 304 and 316L steels currently on the market. Super ductile steel facilitates the forming process and improves the safety of mechanical parts, making it an excellent substitute for Fe-Mn, 304, 316L steels and certain superplastic alloys.

[0096] The foregoing description of the present invention has been presented for purposes of illustration and description only and is not intended to be exhaustive or to limit the invention to the precise form disclosed. Many modifications and variations will be apparent to those skilled in the art.

[0097] The embodiments were chosen and described in order to best explain the principles of the invention and its practical application, 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.

[0098] definition

[0099] It should be noted that the above embodiments are only used to illustrate the technical solution of the present invention and are not intended to be limiting. Although the present invention has described the preferred embodiments in detail, it should be understood by professionals familiar with the art that the embodiments may be modified or replaced by equivalents without departing from the spirit and scope of the technical solution of the present invention. These modifications or equivalents should also be included in the scope of the claims of the present invention.

[0100] As used herein, the terms "approximately", "substantially", "essentially" and "about" are used to describe and explain a small change. When used in conjunction with an event or situation, the term may refer to situations where the event or situation occurs exactly, as well as situations where the event or situation occurs approximately. As used herein, with respect to a given value or range, the term "approximately" generally means within the range of ±10%, ±5%, ±1% or ±0.5% of the given value or range. The range may be indicated herein as from one end point to another or between two end points. Unless otherwise specified, all ranges disclosed in the present invention include endpoints. The term "substantially coplanar" may refer to two surfaces within a few microns located along the same plane, for example, surfaces within a micron, 5 microns, 1 micron or 0.5 micron located along the same face. When referring to "substantially" the same numerical value or characteristic, the term may refer to values ​​within ±10%, ±5%, ±1% or ±0.5% of the average value of the index value.

[0101] Throughout this specification, unless the context requires otherwise, the word "comprise" or variations thereof (such as "comprises" or "comprising") should be understood to imply the inclusion of the stated integer or groups of integers, but not the exclusion of any other integer or groups of integers. It should also be noted that in this disclosure and particularly in the claims and / or paragraphs, terms such as "comprises", "comprised", "comprising" and the like may have the meaning ascribed to them in U.S. patent law, such that they allow for elements not expressly recited, but exclude elements found in the prior art or elements that affect the basic or novel characteristics of the invention.

[0102] Furthermore, throughout the specification and claims, unless otherwise required, the term "include" or variations thereof such as "includes" or "including", will be understood to imply the inclusion of a stated integer or group of integers but not the exclusion of any other integer or group of integers.

[0103] When the description includes the terms "one embodiment", "embodiment", "example embodiment", etc., it means that the described embodiment may include specific features, structures or characteristics, but each embodiment may not necessarily include the specific features, structures or characteristics, and these terms do not necessarily refer to the same embodiment. In addition, when a specific feature, structure or characteristic is described in conjunction with an embodiment, it is considered that it is within the knowledge of those skilled in the art that these features, structures or characteristics can be combined with other embodiments, whether or not explicitly described.

[0104] In the preparation methods described herein, the steps may be performed in any order without departing from the principles of the invention unless a time sequence or an operational order is explicitly stated. In a claim, it should be stated that a step is performed first and then several other steps are performed, which means that the first step is performed before any other steps, but the other steps may be performed in any appropriate order unless the sequence is further listed in the other steps. For example, a claim element listing "step A, step B, step C, step D, and step E" should be understood as step A being performed first and step E being performed last, and steps B, C, and D may be performed in any order between steps A and E, and the order is still within the literal scope of the claim process. A given step or subset of steps may also be repeated. In addition, the specified steps may be performed simultaneously unless the explicit claim language specifies that they are performed separately.

[0105] The term "superductility" or "superplasticity" refers to materials or alloys that exhibit extremely high ductility or deformability under mechanical load. It describes the ability of a material to undergo large plastic deformations without fracture or failure, usually characterized by significant uniform elongation and high total elongation before fracture.

[0106] The term "phase composition" refers to the relative proportions and types of different phases present in a material or alloy. In the case of carbon-doped alloy steel, phase composition refers to the distribution and amount of different phases (e.g., martensite and austenite) within the alloy structure.

[0107] The term "austenite" refers to a specific crystal structure phase in iron and its alloys. It is characterized by a face-centered cubic lattice with closely packed atoms. Austenite has good plasticity and ductility and is a common phase in some steels. Austenite can be formed by cooling high-temperature iron alloys or during solid-state phase transformations, such as the transformation of austenite to a ferromagnetic phase called "martensite". Austenite plays an important role in many metal processing and material applications.

[0108] The term "Kurdjumov-Sachs orientation relationship" or "Nishiyama-Wasserman orientation relationship" refers to a specific crystallographic relationship between two phases in a material. It describes the orientation relationship between the parent phase (usually the high-temperature phase) and the transformed phase (usually the lower-temperature phase) during a phase transition. In the case of the Kurdjumov-Sachs orientation relationship, the crystal planes and directions of the parent and transformed phases are related by a specific mathematical relationship, allowing a coherent interface between the two phases.

[0109] Other definitions of selected terms used herein can be found in the detailed description of the entire invention and application. Unless otherwise defined, all other technical terms used herein have the same meaning as commonly understood by those of ordinary skill in the art to which the invention belongs.

[0110] Those skilled in the art will appreciate that, in light of these teachings, alternative embodiments may be implemented without undue experimentation or deviation without departing from the spirit or scope of the invention as described in the appended claims. The present invention is limited solely by the following claims, which, when considered in conjunction with the above description and drawings, encompass all such embodiments and modifications.

[0111] References: The disclosures of the following references are incorporated herein by reference [1] J. Han, S.-H. Kang, S.-J. Lee, M. Kawasaki, H.-J. Lee, D. Ponge, D. Raabe, Y.-K. Lee, Superplasticity in a lean Fe-Mn-Al steel, Nature Communications 8(1) (2017) 751. [2] TG Langdon, A unified approach to grain boundary sliding in creep and superplasticity, Acta Metall. Mater., 42(7) (1994), 2437-2443. [3] AH Chokshi, AK Mukherjee, TG Langdon, Superplasticity in advanced materials, Materials Science and Engineering: R: Reports 10(6) (1993) 237-274. [4] K. Edalati, T. Masuda, M. Arita, M. Furui, X. Sauvage, Z. Horita, R. Z. Valiev, Room-Temperature Superplasticity in an Ultrafime-Grained Magnesium Alloy, Sci. Rep. 7(1)(2017)2662. [5] HW Swift, Plastic instability under plane stress, Journal of the Mechanics and Physics of Solids, 1(1)(1952) 1-18. [6] E. Nes, Modelling of workhardening and stress saturation in FCC metals, Prog. Mater. Sci. 41(3) (1997) 129-193. [7] B. Gludovatz, A. Hohenwarter, D. Catoor, EH Chan, EP George, RORitchie, A fracture-resistant high-entropy alloy for cryogenic applications, J. Chem. Soc. Am., 345(6201)(2014)1153-1158. [8] D. Raabe, C. C. Tasan, H. Springer, M. Bausch, From High-Entropy Alloys to High-Entropy Steels, 86(10) (2015) 1127-1138.

Claims

1. A carbon-doped alloy steel having room temperature super ductility, characterized in that: The carbon-doped alloy steel comprises the following chemical components: 10-30% by weight of nickel (Ni), 0.01-5% by weight of silicon (Si), 0.01-2% by weight of carbon (C) and 60-90% by weight of iron (Fe), wherein due to the incorporation of C into the FeNiSi alloy, the carbon-doped alloy steel exhibits a room temperature uniform elongation of at least 100%. 2 . The carbon-doped alloy steel according to claim 1 , wherein the carbon-doped alloy steel further comprises 0.1-20% by weight of manganese (Mn).

3. The carbon-doped alloy steel according to claim 1, wherein the incorporation of the carbon into the FeNiSi alloy results in a phase composition transformation from a single martensite phase to a single austenite phase.

4. The carbon-doped alloy steel according to claim 1, wherein the yield strength of the carbon-doped alloy steel after cold rolling is at least 1200 MPa.

5. The carbon-doped alloy steel according to claim 1, wherein the carbon-doped alloy steel elongates uniformly during the entire tensile test without obvious necking.

6. The carbon-doped alloy steel according to claim 1, wherein the carbon-doped alloy steel exhibits a face-centered cubic (fcc) structure with a fully recrystallized morphology.

7. The carbon-doped alloy steel according to claim 1, wherein the tensile strength of the carbon-doped alloy steel is increased by at least 30% compared to its yield strength.

8. The carbon-doped alloy steel according to claim 1, wherein the carbon-doped alloy steel exhibits a true stress in the range of 1500-2200 MPa when the true strain ranges from 80% to 100%.

9. A method for preparing a carbon-doped alloy steel having room temperature super ductility, characterized in that: The method comprises: arc melting one or more raw materials in an argon environment adsorbed by titanium to obtain an ingot; homogenizing the ingot; The ingot is cold rolled at room temperature to reduce the thickness by 40-60%; the cold rolled ingot is recrystallized to obtain carbon-doped alloy steel; the carbon-doped alloy steel is cold rolled to reduce the thickness by 50-80%; and the carbon-doped alloy steel is annealed to adjust the grain size.

10. The method of claim 9, wherein the one or more raw materials include nickel (Ni), silicon (Si), carbon (C), and iron (Fe).

11. The method according to claim 10, wherein the content of the one or more raw materials includes 10-30% by weight of nickel (Ni), 0.01-5% by weight of silicon (Si), 0.01-2% by weight of carbon (C) and 60-90% by weight of iron (Fe).

12. The method of claim 11, wherein the one or more raw materials further comprise 0.1-20% by weight of manganese (Mn).

13. The method according to claim 9, wherein the homogenization reaction temperature is in the range of 1000-1150°C, and the reaction time is in the range of 1-2 hours.

14. The method according to claim 9, wherein the recrystallization reaction temperature is in the range of 1000-1150°C, and the reaction time is in the range of 1-2 hours.

15. The method according to claim 9, wherein the annealing treatment temperature is in the range of 800-1200°C, and the reaction time is in the range of 0.5-30 minutes.

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