Patent Application
20250019809
The present invention generally relates to the field of material science. Specifically, it pertains to the development of a novel alloy steel with exceptional ductility at room temperature. The potential applications of super-ductile alloy steel include, but are not limited to, critical safety components such as front side members, floor side reinforcements, still inners, rear side members, B-pillar reinforcements, roof bows, A-frame reinforcements, etc.
In order to fabricate complex parts with high precision in a single-cycle process, it is crucial to have materials with extremely large ductility, including superplasticity (defined as ultrahigh ductility above 300%)1. However, such ductility is typically achieved at elevated homologous temperatures and often exhibits limited uniform elongation and deformation softening.
Traditionally, achieving such high ductility involves triggering thermally-activated deformation mechanisms like grain-boundary sliding and coble creep at elevated temperatures relative to the melting point (T/Tm)2,3. The prepared alloys are not suitable for mass production due to their excessive energy consumption. Furthermore, the thermally-activated deformation mechanisms lead to deformation softening during super-plasticity and insufficient strength for structural applications1,4. Therefore, it is imperative to develop alloys with low-temperature (especially room-temperature) ultrahigh ductility and good strength, in order to reduce carbon footprints and energy consumption.
The uniform elongation of metallic materials at low homologous temperatures depends on their ability to undergo sufficient strain-hardening, which enables them to keep up with the increasing flow stress5. However, the strain-hardening capability decreases as strain increases, primarily due to the progressive accumulation of dislocations, thus setting the fundamental limit for the room-temperature ductility of metallic materials6. The main approach to enhance room-temperature ductility has been to reduce the stacking fault energy, promoting twinning-induced plasticity (TWIP) or transformation-induced plasticity (TRIP)7. Nevertheless, despite the inclusion of TWIP and TRIP, the room-temperature uniform elongation of metallic alloys has rarely exceeded 100%8.
Accordingly, the objective of the present invention is to provide a carbon-doped alloy steel with super room-temperature ductility. In a first aspect, the carbon-doped alloy steel includes a chemical composition including 10-30 wt. % of nickel (Ni), 0.01-5 wt. % of silicon (Si), 0.01-2 wt. % of carbon (C), and 60-90 wt. % of iron (Fe). The incorporation of carbon into the FeNiSi alloy leads to a transformation in the phase composition, shifting from a single martensite phase to a single austenite phase.
In accordance with one embodiment, optionally the carbon-doped alloy steel further comprises 0.1-20 wt. % of manganese (Mn).
In accordance with one embodiment, the carbon-doped alloy steel has a uniform elongation of at least 100%.
In accordance with one embodiment, the carbon-doped alloy steel after cold rolling has a yield strength of at least 1200 MPa.
In accordance with one embodiment, the carbon-doped alloy steel is uniformly elongated during the whole tensile testing without obvious necking.
In accordance with one embodiment, the carbon-doped alloy steel shows a single face-centered cubic (fcc) structure with a fully recrystallized morphology.
In accordance with one embodiment, the carbon-doped alloy steel has an increased engineering strength of at least 30%.
In accordance with one embodiment, the carbon-doped alloy steel exhibits a true stress in a range of 1500-2200 MPa at a true strain ranging from 80-100%.
The present invention first observes super room-temperature ductility in the FeNiSiC alloy. This remarkable ductility is achieved through sustainable work-hardening capability. Specifically, carbon plays a crucial role in stabilizing the austenite phase, undergoes dynamic redistribution in the FeNiSiC alloy during deformation. This redistribution leads to local variations in stacking fault energy and austenite stability, triggering the sequential activation of multiple deformation mechanisms, i.e., dislocation slipping, deformation twinning, and martensite transformation in the single fcc-phase FeNiSiC alloy. These internal defects progressively refine the structure of the FeNiSiC alloy from tens of micrometers to the nanoscale, resulting in a sustainable work-hardening capability. These findings pave the way for enhancing the ductility of metallic materials by integrating various work-hardening mechanisms through the dynamic redistribution of critical solute elements.
In a second aspect, the present invention provides a method for preparing the carbon-doped alloy steel with super room-temperature ductility, including:
In accordance with one embodiment, the homogenization reaction temperature is in a range of 1000-1150° C., and the reaction time is in a range of 1-2 hours.
In accordance with one embodiment, the recrystallization reaction temperature is in a range of 1000-1150° C., and the reaction time is in a range of 1-2 hours.
In accordance with one embodiment, the annealing reaction temperature is in a range of 800-1200° C., and the reaction time is in a range of 0.5-30 minutes.
The present invention has the following advantages:
Embodiments of the invention are described in more details hereinafter with reference to the drawings, in which:
Steel makes up around 65% of an average automobile's weight. In 2020, the high-strength steel consumption of the auto industry in China reaches 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 sacrificing their properties. Additionally, complex mechanical parts typically require separate fabrication and subsequent assembly, resulting in high energy costs.
In the steel, adding Ni, C, and Si could effectively stabilize the austenite and decrease the martensite transformation start temperature. However, excessive Mn, Ni, Si, and C make the austenite too stable and suppress the occurrence of deformation nanotwins and strain-induced martensite. Adding Si increases the strength of austenite and enhances the resistance to oxidation and corrosion. However, excessive Si facilitates the formation of the brittle phase, decreasing the ductility, toughness, and processability of the steel.
The present invention founds that carbon (C) plays an important role in controlling the deformation mechanisms by dynamically influencing stacking fault energy and austenite stability. The super-ductile carbon-doped steel makes it possible to fabricate parts with complex shapes in one step. In particular, the present invention provides a carbon-doped alloy steel with super room-temperature ductility, the carbon-doped alloy steel includes a chemical composition including 10-30 wt. % of nickel (Ni), 0.01-5 wt. % of silicon (Si), 0.01-2 wt. % of carbon (C), and 60-90 wt. % of iron (Fe). The carbon-doped alloy steel exhibits a room-temperature uniform elongation of at least 100% attributed to the incorporation of C into a FeNiSi alloy.
The basic principle for designing steels with super-ductile involves two aspects: (1) controlling the stability of the austenite phase to ensure that the steel consists primarily of austenite after quenching, and (2) adjusting the stacking fault energy and stability of the austenite to facilitate the sequential activation of various deformation mechanisms, in terms of dislocation, strain-induced twinning, and martensite transformation.
In one embodiment, the carbon-doped alloy steel includes four elements: 10-30 wt. % of nickel (Ni), 0.1-5 wt. % of silicon (Si), 0.01-2 wt. % of carbon (C), and 60-90 wt. % of iron (Fc). “wt %” in this context refers to weight percent, which is a measure of the relative weight of a particular element or component in the alloy steel compared to the total weight of the alloy steel itself.
Preferably, the carbon-doped alloy steel includes 23-25 wt. % of Ni, 0.8-3 wt. % of Si, 0.1-0.5 wt. % of C, and 71.5-76.1 wt. % of Fc.
In another embodiment, the carbon-doped alloy steel includes five elements: 10-30 wt. % of nickel (Ni), 0.1-20 wt. % of manganese (Mn), 0.01-5 wt. % of silicon (Si), 0.01-2 wt. % of carbon (C), and 43-89.88 wt. % of iron (Fc).
Preferably, the carbon-doped alloy steel includes 18-26 wt. % of Ni, 1-10 wt. % of Mn. 0.05-4 wt. % of Si, 0.05-0.9 wt. % of C and 59.1-80 wt. % of Fc.
In one embodiment, the carbon-doped alloy steel exhibits a uniform elongation of at least 100%. For instance, 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%.
In one embodiment, the carbon-doped alloy steel after cold-rolling has a yield strength of at least 1200 MPa. For example, the yield strength is at least 1300 MPa, at least 1350 MPa, at least 1400 MPa after cold-rolling, which is three times higher than the yield strength before cold-rolling.
Preferably, after cold-rolling with a thickness reduction of 76%, the carbon-doped alloy steel could achieve a favorable combination of strength and ductility, specifically a yield strength up to 1400 MPa and a uniform elongation of 40%, surpassing that of most existing Fe—Mn steel, 304, and 316L stainless steels.
In one embodiment, the carbon-doped alloy steel has an increased engineering strength of at least 30%. For instance, the engineering strength increases at least 40%.
These exceptional forming capabilities and distinctive mechanical properties enable the super-ductile steels to be utilized in near-net shape-forming processes. In such processes, the super-ductile steels can be shaped into complex parts while simultaneously attaining sufficient strength for practical applications.
In another aspect, the present invention also provides a method for preparing the carbon-doped alloy steel with super room-temperature ductility. The processing of the super-ductile alloy includes four steps: (1) The alloy was fabricated by arc melting high-purity (99.9 wt. %) raw materials under a Ti-gettered argon atmosphere. (2) To eliminate the possible casting defects, the as-cast ingots were homogenized at 1000-1150° C. for 2 hours followed by cold rolling at room temperature to a thickness reduction of 40-60%, then recrystallized at 1000-1150° C. for 1-2 hours. This step could also be replaced by hot forging and hot rolling at the temperature above the austenitizing temperature. (3) Adjusting grain size by cold-rolling and subsequent recrystallization annealing treatment. The sample was cold rolled with a thickness reduction of 50-80%, followed by an annealing treatment at 800-1200° C. for 0.5-30 mins. The annealing treatment parameters should be chosen according to the thickness of cold-rolled samples.
Given the excellent combination of strength and ductility exhibited by the super-ductile alloy steel after forming, it has the potential to replace many existing advanced high-strength steels in applications such as automobiles, buildings, and shipbuilding. The super-ductile alloy steel is expected to attract considerable attention from automakers and steel enterprises.
The room-temperature super-ductility was realized by doping carbon in Fe—Ni—Si alloys with poor ductility. In one embodiment, the phase composition of super-ductile alloy steel includes 22-24 wt. % of Ni, 1-2.5 wt. % of Si, 0.3-0.5 wt. % of C and 73-76.7 wt. % of Fc. The materials were then homogenized at 1050-1100° C. for 2 h followed by cold rolling at room temperature with a thickness reduction of 60% and then recrystallized at 1050-1100° C. for 2 h. After being cold rolled at room temperature with a thickness reduction of 76% followed by annealing at 1100° C. for 1 min. The prepared super-ductile alloy steel (FeNiSiC alloy) can exhibit a super high uniform elongation of about 174%.
In this example, FeNiSiC alloy was prepared and processed according to the description in the Example 1.
The appearance of the tensile samples (
The significant difference in tensile properties between the FeNiSi and FeNiSiC alloys raises the question of the role of carbon in altering the initial microstructure, thereby transforming the brittle alloy into a quasi-superplastic alloy. To address this issue, a systematic investigation of the pre-deformed microstructures was conducted using techniques such as electron backscatter diffraction (EBSD), X-ray diffraction (XRD), atom probe tomography (APT), and transmission electron microscopy (TEM).
Inverse pole figure (IPF) map is a graphical tool commonly used in materials science and metallurgy to visualize the orientation of crystals and the spatial distribution of grain arrangements. Referring to
The elemental distribution of the two alloys was characterized by energy dispersive spectroscopy (EDS) on TEM. The results revealed that all the elements were uniformly distributed in both FeNiSiC alloys (
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 uniformity of the alloy without carbon segregation in the grain interior. Referring to
Different from the super-ductility at high temperatures that originates from a high strain rate sensitivity, the giant uniform elongation of FeNiSiC alloy stems from a sustainable work-hardening capability. Referring to
Systematically microstructural characterizations reveal that the multiple work-hardening behaviors in the 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 were the main carriers of plastic strains in stage I.
With further straining, deformation twins were activated in the FeNiSiC alloy during stage II.
For example, at εT=45%, high-density nanoscale deformation twins (with an average twin/matrix lamellae of 26±6.4 nm) accompanied by stacking faults were observed, constituting a volume fraction of 30±5% (estimated by TEM). The rotation observed in the selected area electron diffraction pattern indicated that twin boundaries were distorted as a result of their strong interaction with dislocations. The presence of twin boundaries impeded further dislocation slip, thereby enhancing the work-hardening capability of the material. This allows the material to sustain more plasticity without undergoing necking.
Subsequent to deformation twinning, α′-martensite transformation was also evoked to sustain plastic strains in stage III. The results showed that the α′-martensite tended to nucleate in the twinned region, either at the twin boundaries following the Kurdjumov-Sachs orientation relationship (K—S, (111)γ//(1
In addition, the volume fraction of martensite also increased with increasing strain, from 20 vol. % at a true strain of 72% to 34 vol. % at failure. Meanwhile, the volume fraction of deformation twins increased to 80 vol. % when failed. The extensive deformation twinning and the strain-induced α′-martensite transformation in stage III generated a nano-scale martensite-austenite dual-phase structure (
The sequential activation of deformation twinning and α′-martensite transformation is abnormal in the single-phase fcc steel with a homogeneous recrystallized structure. Especially, the FeNiSiC alloy has a stacking fault energy of 22˜27 mJ, which is supposed to be a TWIP alloy. Usually, a hierarchical structure is necessary to unify the two work-hardening mechanisms in metastable alloys, in which the grain-size hierarchy induces a wide variation in phase stability and thereby encourages the concurrent of different deformation mechanisms.
Chemical distribution plays a vital role in determining the stability of austenite. To investigate the microstructural origin of the multiple deformation mechanisms, the present invention characterized the evolution of chemical distribution in the FeNiSiC alloy during tensile tests using the correlative APT-TEM analysis.
Unexpectedly, the results showed that carbon, one of the important austenite stabilizers, tended to dynamically redistribute during the tensile testing at room temperature (
TWIP and α′-martensite transformation induced plasticity has been extensively studied in many alloys, which, however, could only push the room-temperature uniform elongation to a maximum of 103% and 50% (
The pronounced TWIP and TRIP effects in the FeNiSiC alloy may be attributed to the following factors. Firstly, the density of deformation twins and martensite is super high, dividing the coarse grains into nanostructures, i.e., nanotwins with an average thickness of 25±6.4 nm and nanoscale martensite/austenite dual-phase structure with an average size of martensite of 13±0.9 nm (
In summary, the present super-ductile steels emerge as excellent candidates for near-net shape-forming processes. For instance, when subjected to cold-rolling 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 surpass those of most Fe—Mn, 304, and 316L steels currently available. Super-ductile steels facilitate the forming process and enhance the safety of mechanical parts, making them superior alternatives to Fe—Mn, 304, 316L steels, as well as certain superplastic alloys.
The foregoing description of the present invention has been provided for the purposes of illustration and description. It is not intended to be exhaustive or to limit the invention to the precise forms disclosed. Many modifications and variations will be apparent to the practitioner skilled in the art.
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 that are suited to the particular use contemplated.
It should be noted that the above-mentioned embodiments are only for illustrating the technical solutions of the present invention and not for limiting, and although the present invention has been described in detail with reference to the preferred embodiments, it should be understood by those skilled in the art that modifications or equivalent substitutions may be made on the technical solutions of the present invention without departing from the spirit and scope of the technical solutions of the present invention, which should be covered by the claims of the present invention.
As used herein, terms “approximately”, “basically”, “substantially”, and “about” are used for describing and explaining a small variation. When being used in combination with an event or circumstance, the term may refer to a case in which the event or circumstance occurs precisely, and a case in which the event or circumstance occurs approximately. As used herein with respect to a given value or range, the term “about” generally means in the range of ±10%, ±5%, ±1%, or ±0.5% of the given value or range. The range may be indicated herein as from one endpoint to another endpoint or between two endpoints. Unless otherwise specified, all the ranges disclosed in the present disclosure include endpoints. The term “substantially coplanar” may refer to two surfaces within a few micrometers (μm) positioned along the same plane, for example, within 10 μm, within 5 μm, within 1 μm, or within 0.5 μm located along the same plane. When reference is made to “substantially” the same numerical value or characteristic, the term may refer to a value within ±10%, ±5%, ±1%, or ±0.5% of the average of the values.
Throughout this specification, unless the context requires otherwise, the word “comprise” or variations such as “comprises” or “comprising”, 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. It is also noted that in this disclosure and particularly in the claims and/or paragraphs, terms such as “comprises”, “comprised”, “comprising” and the like can have the meaning attributed to it in U.S. Patent law; e.g., they allow for elements not explicitly recited, but exclude elements that are found in the prior art or that affect a basic or novel characteristic of the present invention.
Furthermore, throughout the specification and claims, unless the context requires otherwise, the word “include” or variations 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.
References in the specification to “one embodiment”, “an embodiment”, “an example embodiment”, etc., indicate that the embodiment described can include a particular feature, structure, or characteristic, but every embodiment may not necessarily include the particular feature, structure, or characteristic. Moreover, such phrases are not necessarily referring to the same embodiment. Further, when a particular feature, structure, or characteristic is described in connection with an embodiment, it is submitted that it is within the knowledge of one skilled in the art to affect such feature, structure, or characteristic in connection with other embodiments whether or not explicitly described.
In the methods of preparation described herein, the steps can be carried out in any order without departing from the principles of the invention, except when a temporal or operational sequence is explicitly recited. Recitation in a claim to the effect that first a step is performed, and then several other steps are subsequently performed, shall be taken to mean that the first step is performed before any of the other steps, but the other steps can be performed in any suitable sequence, unless a sequence is further recited within the other steps. For example, claim elements that recite “Step A, Step B, Step C, Step D, and Step E” shall be construed to mean step A is carried out first, step E is carried out last, and steps B, C, and D can be carried out in any sequence between steps A and E, and that the sequence still falls within the literal scope of the claimed process. A given step or sub-set of steps can also be repeated. Furthermore, specified steps can be carried out concurrently unless explicit claim language recites that they be carried out separately.
The term “super-ductile” or “superplasticity” refers to a material or alloy that exhibits an exceptionally high level of ductility or deformability under mechanical loading. It describes a material's ability to undergo large plastic deformation without fracture or failure, typically characterized by a significant uniform elongation and high total elongation before fracture.
The term “phase composition” refers to the relative proportions and types of different phases present in a material or alloy. In the context of the carbon-doped alloy steel, the phase composition refers to the distribution and amounts of different phases, such as martensite and austenite, within the alloy structure.
The term “austenite” is a specific crystal structure phase in iron and its alloys. It is characterized by a face-centered cubic lattice with close-packed atomic arrangement. Austenite exhibits good plasticity and ductility and is a common phase in certain steels. Austenite can be formed by cooling high-temperature iron alloys or during solid-state phase transformations, such as the transformation of austenite to the ferromagnetic phase called “martensite”. Austenite plays an important role in many metal processes and material applications.
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 a high-temperature phase) and the transformed phase (usually a lower-temperature phase) during a phase transformation. In the case of Kurdjumov-Sachs orientation relationship, the crystal planes and directions of the parent and transformed phases are related by specific mathematical relationships, allowing for a coherent interface between the two phases.
Other definitions for selected terms used herein may be found within the detailed description of the present invention and apply throughout. Unless otherwise defined, all other technical terms used herein have the same meaning as commonly understood to one of ordinary skill in the art to which the present invention belongs.
It will be appreciated by those skilled in the art, in view of these teachings, that alternative embodiments may be implemented without undue experimentation or deviation from the spirit or scope of the invention, as set forth in the appended claims. This invention is to be limited only by the following claims, which include all such embodiments and modifications when viewed in conjunction with the above specification and accompanying drawings.
