Superelastic and Shape-Memory Fe-Based Wire and Direct Wire Casting Thereof

BACKGROUND

This invention relates generally to ferrous superelastic and shape-memory alloys, and more particularly relates to ferrous superelastic and shape-memory alloy wire and similar structures.

Shape-memory alloys (SMAs) are a unique class of functional materials that can exhibit two notable properties, namely, the shape-memory effect, which refers to the material's ability to recover a memorized shape upon heating, and superelasticity, which allows the material to fully recover large strains on the order of 10% by simply removing the applied load. These properties arise owing to the ability of SMAs to undergo a thermoelastic martensitic phase transformation, which permits solid-state transduction between heat and strain, and naturally gives rise to a variety of actuation, energy-harvesting, and energy-damping applications for SMA materials.

The most well-developed SMAs are NiTi-based SMAs, which have been developed for specialized commercial applications such as orthodontics, glasses frames, and medical stents. While such applications have been well-addressed by NiTi SMAs, NiTi imposes a high cost on a metals basis and the resulting SMA structures can exhibit relatively low ductility, e.g., <30% plastic strain before fracture, unless numerous, costly, intermediate annealing steps are undertaken during shape forming of NiTi; these factors have limited more widespread application of NiTi SMAs. As a result, there is sustained interest in alternative SMAs, such as Cu-based and Fe-based, that is, ferrous, SMAs, which have significantly lower material costs. Fe-based SMAs have the additional advantage that they can be produced with conventional low-cost ferrous alloy processing.

Many Fe-based alloys and ferrous steels exhibit a martensitic transformation between a face-centered cubic (fcc) γ austenite phase and a body-centered tetragonal (bct) α′ martensite phase, and such has been exploited in the production of high-strength martensitic steels. The martensitic transformation in ferrous materials is not usually thermoelastic, however, because this transformation generally causes considerable plastic deformation, disabling the ability of a ferrous structure to return to a pre-martensitic-transformation shape upon reversion. Thus, shape-memory and superelasticity are generally not observed in ferrous materials such as Fe-based alloys.

It has been shown that when a distribution of precipitates is imposed in the austenite matrix of a ferrous alloy, the martensitic transformation can become thermoelastic. A thermal ageing process has been proposed for producing such a precipitate distribution in an austenite matrix. But even with the inclusion of a precipitate population, only tiny superelastic strains, of no more than about 0.7%, have been achievable; no precipitate population distribution in a ferrous alloy has heretofore enabled reasonable superelastic strain performance in the ferrous alloy.

In an effort to overcome this limitation in Fe-based alloys, it has been proposed to impart strong crystallographic texturing of the grains in a polycrystalline Fe-based alloy. Crystallographic texturing of a polycrystalline material generally refers to processing of the material in a manner that increases the alignment of the grains within the material. It has been proposed to employ cold working processes, such as cold-rolling, and subsequent recrystallization to produce crystallographic texturing of polycrystalline Fe-based alloys. Such crystallographic texturing increases grain alignment, possibly reducing crystallographic incompatibilities during martensitic transformation, and may preferentially orient grains in such a way that the grains can sustain relatively larger transformation strains along an axis of interest.

Conventional Fe-based alloy structures having random or weakly textured polycrystalline grains exhibit very limited ductility and essentially no transformation strain. The enhanced martensitic transformation compatibility achieved by strong texturing is understood to be critically required to achieve ductility and superelasticity in polycrystalline ferrous alloys. But cold working processes for strongly texturing polycrystals add significant cost and complexity to Fe-based material production. As a result, progress towards the commercial viability of ferrous SMAs has been significantly limited, with thermoelastic properties generally considered unattainable in Fe-based alloy structures in practice.

SUMMARY

Herein is provided a ferrous shape memory alloy (SMA) wire and processes for production of ferrous shape memory alloy wire that do not require crystallographic texturing processes to achieve superior superelastic and SMA wire properties. In one embodiment, there is provided a shape memory alloy wire that includes an elongated wire body with a longitudinal-axis length of iron alloy material and having a cross-sectional wire diameter that is less than about 1 millimeter. The iron alloy material has an oligocrystalline crystallographic morphology along the longitudinal-axis length. The iron alloy material has a γ-fcc crystallographic matrix and a volume fraction of γ′-L1₂crystallographic precipitates in the γ-fcc crystallographic matrix.

In a method provided herein for producing the ferrous SMA wire, there is mixed a shape memory alloy composition consisting of between about 20 at. % and about 35 at. % nickel (Ni), no more than about 20 at. % cobalt (Co), between about 7 at. % and about 20 at. % aluminum (Al), no more than about 12 at. % total of an element selected from tantalum (Ta), niobium (Nb), titanium (Ti), tungsten (W), molybdenum (Mo), vanadium (V), silicon (Si), chromium (Cr), copper (Cu), antimony (Sb), tin (Sn), zirconium (Zr), hafnium (Hf), thallium (Tl), germanium (Ge), gallium (Ga), and indium (In), and no more than about 2 at. % total of an element selected from boron (B), carbon (C), and nitrogen (N), with iron (Fe) as a remaining balance of the iron alloy composition, to obtain a resulting shape memory alloy mixture. The shape memory alloy mixture is heated to form a melted shape memory alloy, and then the melted shape memory alloy is cast to form a polycrystalline ferrous shape memory alloy wire having a length of at least about 1 meter and a diameter of no more than about 300 microns. The polycrystalline ferrous shape memory alloy wire is then solutionized for a duration that transforms the polycrystalline ferrous shape memory alloy wire to an oligocrystalline ferrous shape memory alloy wire and that produces a γ-fcc phase matrix in the oligocrystalline ferrous shape memory alloy wire. The solutionized oligocrystalline ferrous shape memory alloy wire is then aged to produce in the oligocrystalline ferrous shape memory alloy wire a volume fraction of at least about 10% of γ′-L1₂precipitates. Each of the casting, solutionizing, and ageing steps, and any processing of the solutionized and aged oligocrystalline ferrous shape memory alloy wire subsequent to the ageing step, is conducted without solid state deformation of the ferrous shape memory alloy wire that increases crystallographic texture of the oligocrystalline ferrous shape memory alloy wire.

This method for producing the oligocrystalline ferrous SMA wire provided herein enables rapid, continuous wire-casting for production of kilometer lengths of continuous wire. The high-throughput, low-cost processing methodology of this method, combined with the substantially superior shape-memory properties of the oligocrystalline ferrous SMA wire thereby produced, enables widespread commercial application of ferrous shape memory alloys, and ferrous shape memory wires in particular. Further features and advantages will be apparent from the following description and accompanying drawings, and from the claims.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1A is a schematic view of a length of oligocrystalline ferrous SMA wire provided herein;

FIG. 1B is a schematic view of a section of the length of oligocrystalline ferrous SMA wire of FIG. 1A, exhibiting the oligocrystalline structure of the wire as-provided herein;

FIG. 2 is a schematic view of a section of an oligocrystalline ferrous SMA wire provided herein, exhibiting a weak crystallographic texturing of the grains of the oligocrystalline structure wire structure;

FIG. 3A is a flow chart of the steps of a method provided herein for producing the oligocrystalline ferrous SMA wire provided herein;

FIGS. 3B-3C are each a schematic view of an example wire casting arrangement for producing the oligocrystalline ferrous SMA wire provided herein;

FIG. 4 is a schematic view exhibiting particulars of the wire casting arrangement of FIG. 3B;

FIG. 5 is an Inverse Pole Figure (IPF) showing the measured distribution of the orientation of grains of an experimental oligocrystalline ferrous SMA wire as-provided herein after wire casting;

FIG. 6 is a 3-D reconstruction of a solutionized and aged experimental oligocrystalline ferrous wire as-provided herein, the reconstruction produced from Atom Probe Tomography (APT) to exhibit of γ′-L1₂precipitates of the experimental wire;

FIG. 7 is an Inverse Pole Figure (IPF) showing the measured distribution of the orientation of grains of an experimental oligocrystalline ferrous SMA wire as-provided herein after solutionizing and ageing the experimental wire;

FIG. 8 is an image produced by Electron Backscatter Diffraction (EBSD) of an experimental oligocrystalline ferrous SMA wire as-provided herein after solutionizing and ageing the experimental wire, exhibiting the oligocrystalline bamboo grain structure of the experimental wire;

FIG. 9A is a plot of measured engineering stress as a function of engineering strain for an experimental oligocrystalline ferrous SMA wire as-provided herein after solutionizing and ageing the experimental wire, over 38 superelastic cycles with a 700 MPa load;

FIG. 9B is a plot of measured engineering stress as a function of engineering strain for an experimental oligocrystalline ferrous SMA wire as-provided herein after solutionizing and ageing the experimental wire, over 20 superelastic cycles with a 1 GPa load;

FIG. 10A is a plot of measured engineering stress as a function of engineering strain for an experimental oligocrystalline ferrous SMA wire as-provided herein after solutionizing and ageing the experimental wire, for each of five different temperatures between −40° C. and 30° C.;

FIG. 10B is a plot of the measured temperature dependence of martensite start (σM_s) and austenite finish (σ_Af) stresses, measured to be 5.4 MPa ° C.⁻¹, for an experimental oligocrystalline ferrous SMA wire as-provided herein after solutionizing and ageing the experimental wire; and

FIG. 11 is a plot of measured engineering stress as a function of engineering strain for an experimental oligocrystalline ferrous SMA wire as-provided herein after solutionizing and ageing the experimental wire, at a temperature of −100° C.

DETAILED DESCRIPTION

It is recognized by the inventors herein that fundamentally, to fully establish shape-memory properties in a polycrystalline ferrous alloy, the crystallographic arrangement of a polycrystalline ferrous alloy must operate to reduce grain constraints and intergranular precipitation during martensitic transformation. But the inventors herein have discovered that solid state deformation, i.e., mechanical processing to deform, a polycrystalline ferrous alloy to impose strong crystallographic texturing, for reducing grain constraints and intergranular precipitation, is not required, and indeed, can be eliminated, in production of a ferrous alloy while at the same time achieving superior shape-memory properties that are comparable to that of single crystal ferrous alloy materials.

Referring to FIG. 1A, there is schematically shown an example wire 10 in one embodiment of a ferrous SMA structure provided herein. The wire 10, like all ferrous SMA structures provided herein, is an oligocrystalline structure of a ferrous alloy and includes a dispersed population of γ′ phase precipitates in a γ phase matrix. In various embodiments provided herein, the oligocrystalline structure is a 1-dimensional structure. In embodiments herein, the oligocrystalline structure is configured as, for example, a wire, tape, pillar, fiber, ribbon, tube, sheet, cantilever, wall, bridge, open cell foam, or other structure having one dimension that is substantially larger than the other two dimensions of the structure. In other words, in various embodiments, the structures provided herein include a feature having an extent that is substantially larger than the extent of the other two dimensions of the structure. Examples of such features provided in embodiments herein include the length of a wire, tape, pillar, fiber, or sheet.

FIG. 1B is a schematic view of a section 12 of the oligocrystalline ferrous SMA wire 10 of FIG. 1A. As evidenced by the wire section of FIG. 1B, the wire 10, like all ferrous SMA structures provided herein, is an oligocrystalline structure. In one embodiment, the term “oligocrystalline structure” as used herein refers to a structure having a polycrystalline morphology in which the average size of the grains in the material composing the structure is greater than about at least half of the extent of a dimension of the structure. For example, grains 14 of the wire section 12 are shown to fully extend across the diameter, d, of the wire section 12, or to extend about half-way across the diameter of the wire 12. This grain configuration is characteristic of an oligocrystalline structure. For an oligocrystalline wire, the grains are in a bamboo or near-bamboo-like arrangement along the length of the wire. In another embodiment, an oligocrystalline structure provided herein exhibits a total surface area that is greater than the total area of the polycrystalline grain boundaries within the structure, in the manner prescribed in U.S. Pat. No. 8,282,746, issued Oct. 9, 2012; and in the manner prescribed in U.S. Pat. No. 10,214,798, issued Feb. 26, 2019, the entirety of both of which are hereby incorporated by reference.

In various embodiments herein, the oligocrystalline ferrous SMA wire 10 is characterized by a total surface area that is greater than the total area of crystalline grain boundaries within the wire. This condition is characteristic of a substantially full bamboo structure in which single grains span the entire diameter of the wire. With this arrangement, the grains of the wire are coordinated predominantly by unconfined free surfaces rather than by rigid boundaries with other grains within the wire. In contrast with bulk polycrystals in which substantially all grains have zero degrees of freedom and must undergo martensitic transformation in a coordinated fashion with neighboring grains in all directions, very few grains in the oligocrystalline wire 10 are constrained on all sides.

In addition, by requiring grains to either fully span a structural feature or extend at least about half-way across a structural feature, the ferrous oligocrystalline SMA structures provided herein, like the oligocrystalline ferrous SMA wire 10, contain relatively few triple junctions of grains, which are the most severe sites of stress concentration due to incompatible transformation deformation between neighboring grains; the oligocrystalline structures provided herein thereby are optimized to have relatively few source points of incompatibility stresses. Further, as shown in FIG. 1B, most grains in an oligocrystalline ferrous SMA wire provided herein have a region that is exposed to some amount of free surface, enabling internal stresses to be partially relaxed. In embodiments herein, the surface-to-volume ratio of an oligocrystalline structure is maximized, as in an elongated wire structure, to optimize the exposure of grains to a free surface for enabling internal stresses to be relaxed during martensitic transformation. The inventors herein have discovered that with these characteristics, the ferrous oligocrystalline SMA structures herein have sufficient reduction in sources of incompatibility stresses, coupled with enhanced stress relaxation, to reduce the peak internal stress during martensitic transformation to a level that is below the characteristic fracture stress of the structure, and also to enable more complete martensitic transformation and less strain cancellation between martensite variants, thereby achieving a large martensitic transformation strain.

In one embodiment provided herein, the oligocrystalline ferrous SMA wire 10 has a wire diameter that is no more than about 1 mm. In a further embodiment, the oligocrystalline ferrous SMA wire diameter is no more than about 300 μm, and in a further embodiment the diameter is no more than 100 μm; in a further embodiment, the oligocrystalline ferrous SMA wire diameter is between about 50 μm and about 300 μm. In various further embodiments provided herein, the oligocrystalline ferrous SMA wire has a wire diameter that is between about 1 μm and about 1000 μm. The term “diameter” refers herein to the cross-sectional extent of the wire as shown in FIG. 1B. For other structures provided herein, the connotation of the term diameter is extended to refer to the smallest dimension of such, e.g., the thickness of an oligocrystalline ribbon, tape, or sheet, and is not intended to imply that the cross-section of the structure, such as a wire cross-section, must be symmetric in the plane. The wire cross section can be circular, oval, or other geometry, and need not be symmetrical; “diameter” refers to a selected extent in the cross section.

In various embodiments herein, the length of an elongated oligocrystalline structure such as oligocrystalline ferrous SMA wire is at least about 10 cm; in more preferred embodiments, the length is at least about 1 mm. But the methods provided herein for production the oligocrystalline ferrous SMA wire enable continuous wire production of kilometers in length. In one embodiment, the oligocrystalline ferrous SMA wire length is at least about 1 meter. In other various embodiments, the length of the oligocrystalline ferrous SMA wire is greater than about 10 meters; in other embodiments, the length of the oligocrystalline ferrous SMA wire is more than 100 m, and in further embodiments, the length of the oligocrystalline ferrous SMA wire is more than 1 km.

In one embodiment provided herein, the crystallographic texture of the ferrous oligocrystalline structure is not a strong crystallographic texture. In further embodiments herein, the crystallographic texture of the ferrous oligocrystalline structure, such as an oligocrystalline wire, is weak or very weak. In one embodiment, there is an absence of strong crystallographic texturing. Referring to FIG. 2, there is shown a section 25 of an oligocrystalline wire provided herein with very weak crystallographic texture, as represented by the generally random arrowhead directions of the grains. This generally substantially random grain orientation is conventionally understood to result in brittle polycrystalline material that is not superelastic. But the inventors herein have discovered that oligocrystalline structures having weak or no crystallographic texture as provided herein can demonstrate superior shape memory properties when arranged in the oligocrystalline morphology provided herein. Thus, in one embodiment, the ferrous oligocrystalline wire exhibits a substantially random grain orientation.

In contrast, crystallographic texturing, which generally causes grains to align in a common orientation, is not required herein. As a result, solid state processing, that is, mechanical processing, is not required, and in one embodiment herein is prohibited, so that no mechanical processes increase the crystallographic texturing of the ferrous oligocrystalline wire. In other words, it is preferred that no mechanical, solid state processes, such as drawing, extruding, or other process, in the production of the ferrous oligocrystalline wire, increase the crystallographic texturing of the wire.

In practice, the crystallographic texture of a structure can be determined by Electron Backscatter Diffraction (EBSD), by X-ray diffraction, by neutron diffraction, or by other suitable technique. One convenient quantitative measure of crystallographic texture is the value of the so-called “multiples of random distribution,” or mrd, of a polycrystalline material. The mrd of a polycrystalline material measures how many times more likely it is to obtain a given crystallographic grain orientation relative to a random grain orientation distribution. To determine the mrd of a given polycrystalline material, EBSD or other technique is employed to determine the orientation of multiple grains at different sites in the material, and a normalized probability density function can be produced with the identified orientations. Then for each identified orientation, the probability of the orientation's existence, as-expressed by the probability density function, divided by the probability of a uniform probability distribution, expresses the mrd of the polycrystalline material.

In embodiments provided herein, it is required that a ferrous oligocrystalline structure exhibit no strong crystallographic texture, or in other words, exhibit an absence of strong crystallographic texturing. In one such embodiment, the crystallographic texture of a ferrous oligocrystalline structure is no greater than that given for mrd=15; in other words, mrd≤15. In other embodiments herein, it is required that a ferrous oligocrystalline structure exhibit no more than weak crystallographic texture. In one such embodiment, the crystallographic texture of a ferrous oligocrystalline structure is no greater than that given for mrd=10; in other words, mrd≤10. In other embodiments provided herein, it is required that the crystallographic texture of a ferrous oligocrystalline structure exhibit no more than very weak crystallographic texture. In one such embodiment, the crystallographic texture of a ferrous oligocrystalline structure is no greater than that given for mrd=5; in other words, mrd≤5. As explained above, the generally random grain orientation associated with very weak texturing is conventionally understood to result in brittle polycrystalline material that is not superelastic. But the inventors herein have discovered that oligocrystalline structures having weak can demonstrate superior shape memory properties when arranged in the manner provided herein.

In embodiments provided herein, the oligocrystalline morphology of the ferrous SMA structures provided herein includes a γ-phase matrix in which is dispersed a non-zero population of γ′-phase precipitates. The term precipitate is used herein to refer to a material region having a different crystal structure and/or a different chemical composition than the adjacent matrix material. In one embodiment provided herein, the γ-phase matrix has a face-centered cubic (fcc) crystal structure and the γ′-phase precipitates have an L1₂crystal structure. The γ′-L1₂precipitates can be any shape, such as cuboidal or near-spherical, and the precipitate shape can change as a result of changes to chemical composition or heat treatment, e.g., to become more ellipsoidal, spherical, cuboidal, or other shape. In one embodiment, each γ′-phase precipitate has an extent that is no more than about 100 nm; in other embodiments, each γ′-phase precipitate has an extent that is between about 1 nm and about 100 nm, and in further embodiments, each γ′-phase precipitate has an extent that is no more than about 20 nm. The precipitate extent specified here refers to the dominant dimension of a precipitate, e.g., the diameter of a sphere, the major axis of an ellipsoid, the cube diagonal of a cuboid, or other precipitate dimension. Transmission Electron Microscopy (TEM), Atom-Probe Tomography (APT), or other suitable imaging can be employed to identify and measure the precipitates in a polycrystalline material to determine if the precipitates meet the criteria given above.

The γ′-phase precipitates constitute a volume fraction of the oligocrystalline ferrous SMA structure volume. In various embodiments, the population of γ′-phase precipitates constitutes at least about 10% by volume of the structural material, with γ-phase morphology constituting at least a portion of the remaining material volume. In further various embodiments, the population of γ′-phase precipitates constitutes between about 20% and about 50% by volume of the structural material, with γ-phase morphology constituting at least a portion of the remaining material volume. In other embodiments, the population of γ′-phase precipitates constitutes at least about 30% by volume of the structural material, with γ-phase morphology constituting at least a portion of the remaining material volume, and in other embodiments, the population of γ′-phase precipitates constitutes at least about 40% by volume of the structural material, with γ-phase morphology constituting at least a portion of the remaining material volume.

In embodiments herein, the number density of the precipitate population is at least about 1×10²³m⁻³. It can be preferred to provide a generally uniform dispersion of the γ′-phase precipitate population within the γ-phase matrix morphology for good functional fatigue properties and for good property uniformity through the oligocrystalline ferrous SMA structure. As explained in detail below, uniform composition of a starting alloy material, and uniform heating of the alloy material during heat treatments, can substantially enforce a generally uniform dispersion of precipitates within the matrix morphology of the oligocrystalline structure.

The oligocrystalline ferrous SMA structure provided herein includes a ferrous alloy compound, that is, includes iron (Fe) in the composition. In embodiments provided herein, the oligocrystalline material composition includes at least about 40 atomic % (at. %) iron. In other embodiments herein, the oligocrystalline material composition includes no more than about 80% iron. In various embodiments, the oligocrystalline composition is an alloy including between about 20-35 at. % nickel (Ni), no more than about 20 at. % cobalt (Co), between about 7-20 at. % aluminum (Al), no more than about 12 at. % total of an element selected from tantalum (Ta), niobium (Nb), titanium (Ti), tungsten (W), molybdenum (Mo), vanadium (V), silicon (Si), chromium (Cr), copper (Cu), antimony (Sb), tin (Sn), zirconium (Zr), hafnium (Hf), thallium (Tl), germanium (Ge), gallium (Ga), and indium (In), and no more than about 2 at. % total of an element selected from boron (B), carbon (C), and nitrogen (N), with Fe included as the remaining balance of the alloy composition. In embodiments herein, 0 at. %-1 at. % of the element boron is included in the ferrous alloy. In other embodiments herein, 0.03 at. %-0.1 at. % of the element boron is included in the ferrous alloy. In various embodiments, the alloy is Fe-(20-35)Ni-(0-20)Co-(7-15)Al-(0-10)(Ta, Nb, Ti, Mo, V, W)-(0-1)B (atomic %). In other various embodiments, the alloy is Fe-(27-30)Ni-(12-18)Co-(9-13)Al-(1-3)(Ta, Nb, Ti)-(0-0.1)B (atomic %). In other various embodiments, the alloy is Fe-28Ni-17Co-11.5Al-2.5Ta-0.05B (atomic %).

In one embodiment provided herein, the alloy composition of the γ′-phase precipitates includes a lower fraction of Fe than the alloy composition of the γ-phase matrix surrounding the precipitates. In a further embodiment provided herein, the alloy composition of the γ′-phase precipitates includes a lower fraction of Co than the alloy composition of the γ-phase matrix surrounding the precipitates. In other embodiments, the alloy composition of the γ′-phase precipitates includes a higher fraction of Ni than the alloy composition of the γ-phase matrix surrounding the precipitates; the alloy composition of the γ′-phase precipitates also can include a higher fraction of Al than the alloy composition of the γ-phase matrix surrounding the precipitates, and similarly, the alloy of the γ′-phase precipitates can include a higher fraction of Ta than the alloy composition of the γ-phase matrix surrounding the precipitates.

The inventors herein have discovered that unexpectedly, the oligocrystalline ferrous alloy structures provided herein, having an alloy composition as-specified above and including a population of γ′-phase precipitates in an oligocrystalline γ-phase matrix, exhibit superior SMA properties, without the use of solid state deformation to impose strong crystallographic texturing. In one embodiment, oligocrystalline ferrous alloy wire provided herein is superelastic at room temperature; in embodiments herein, the oligocrystalline ferrous alloy wire exhibits superelastic behavior over a temperature range of −40° C. and 30° C.

In further embodiments, the oligocrystalline ferrous alloy wire provided herein exhibits a superelastic strain of at least about 1% after training with a γ′-L1₂precipitate population of at least 10% by volume in the γ-phase matrix; in one embodiment, the 1% strain is achieved with a crystallographic texturing having mrd≤15. In other various embodiments, the oligocrystalline ferrous alloy wire provided herein exhibits a superelastic strain of at least about 5%, and exhibits a tensile strength of at least about 1 GPa after training with a γ′-L1₂precipitate population of at least 25% by volume in the γ-phase matrix. In one embodiment, the 5% strain and 1 GPa tensile strength is exhibited with a crystallographic texturing of mrd≤15. In other embodiments, the oligocrystalline ferrous SMA wire provided herein exhibits a superelastic strain of at least 7% and exhibits a tensile strength of at least 1.2 GPa after training with a γ′-L1₂precipitate population of at least 25% by volume in the γ-phase matrix; in further embodiments, the 7% strain and 1.2 GPa tensile strength is obtained in the absence of strong crystallographic texturing, that is, with mrd≤15.

In addition, the stress hysteresis of the oligocrystalline ferrous alloy wire provided herein is considerably lower than that of conventional textured, polycrystalline, ferrous alloy wire, and is instead more consistent with that observed in single-crystal ferrous alloy wires. This unexpectedly superior performance can be attained without added crystallographic texturing, and therefore eliminates the need for solid state deformation processes, thermomechanical processes, or directional solidification processes to increase crystallographic texturing. As a result, the oligocrystalline ferrous alloy structures provided herein enable widespread commercial application of ferrous SMA alloys with efficient, low-cost production.

Referring to FIG. 3A, in one embodiment, there is provided herein a process 100 for producing oligocrystalline ferrous SMA wire as described above. In the process, described in detail below, there is conducted a first step of wire casting 110 to produce a ferrous alloy wire. In general, any process of wire casting can be employed; wire casting methods including various forms of melt spinning, for example, by casting onto a cooled rotating mandrel, drum, slotted mandrel or spinning disk can be employed. In one embodiment, an In Rotation Water Melt-Spinning (INROWASP) process is employed to cast ferrous alloy wire. In embodiments herein, in general, any casting process can be employed in which a molten jet of ferrous alloy is ejected onto a cooling medium. The cooling medium can be solid, liquid, or gaseous.

The ferrous alloy wire resulting from the wire casting step 110 is then subjected to a step of wire solutionization 115 to achieve a γ-fcc phase morphology and oligocrystalline ferrous alloy structure. The resulting oligocrystalline wire is then subjected to a step of wire thermal ageing 120 to form a dispersion of γ′-L1₂precipitates in the oligocrystalline ferrous alloy wire structure. This concludes the oligocrystalline ferrous wire production.

In one embodiment, the method 100 includes the prohibition 125 of solid state deformation processes that increase crystallographic texturing in the oligocrystalline ferrous wire. In various embodiments, solid state deformation for crystallographic texturing is absent from the method 100. In other various embodiments, solid state, i.e., mechanical deformation, processes specifically for crystallographic texturing are not required, and are preferably not included, in the wire production process; solid state processing such as cold working, thermomechanical processing, directional solidification, or other process, to cause crystallographic texturing of the ferrous alloy wire, are not to be included, as they add time, cost, and complexity to the overall material processing, and are unnecessary to achieve the superior performance characteristics provided herein. Cold working can be employed, for example, as wire drawing after wire casting to fine tune wire diameter, can be employed to control wire surface conditions, or for other consideration. But it is preferred that no solid state processing for increase of crystallographic texturing be conducted.

Turning to embodiments of the method 100 of FIG. 3A, in the step 110 of wire casting, any suitable casting process can be employed. Referring to FIGS. 3B and 3C, an In Rotation Water Melt-Spinning arrangement 30, a melt spinning solid-drum arrangement 31, or other suitable arrangement can be employed. In either arrangement, a ferrous wire 10 is produced, either in a liquid medium 42 or on a solid surface 33, e.g., a solid wheel. In general, a molten alloy is fed under pressure 38 from a heated crucible through an ejection nozzle to form a molten jet that is cooled by exposure to a selected cooling medium, such as a gaseous or liquid medium 42 or solid surface 33. For many commercial applications, the embodiment of FIG. 3C can be preferred, in which a solid cooling surface is provided.

In one example embodiment of a liquid-cooled arrangement, referring now to FIG. 4, there is provided a crucible, such as a cylindrical quartz crucible 32 having a nozzle 34 arranged for output of the wire 10 of FIG. 1A, or for output of ribbon, or other structure there from. The crucible is fixedly positioned, e.g., by a manipulator, above a horizontal face 38 of an open-faced vertical rotating drum wheel 40. The drum wheel includes walls on each side of the horizontal face 38 for holding a quenching/casting medium 42. The drum wheel is rotated, in the direction shown in the figure, in a manner that is controlled for SMA structure formation.

The crucible 32 is arranged adjacent to induction coils 36 or other suitable heating mechanism, for melting SMA material that is provided within the crucible to form wire, ribbon, or other structure at the nozzle 34. The crucible is also connected to a source of pressure 38, such as gas pressure, for controllably forcing, or ejecting, melted alloy material out of the nozzle 34. Other pressure arrangements, as well as crucible heating arrangements, can be employed as-suitable for a given application. In production of SMA structures such as SMA wire, bulk solid pieces of ferrous alloy components are loaded into the crucible; the ferrous alloy compositions given above are preferred. With the bulk solid SMA material loaded in the crucible, the crucible is then evacuated and a selected inert gas, such as argon gas, is continuously flowed through the crucible; the material is then heated, and upon reaching the desired temperature, the pressure can be increased to between about 1 bar and about 10 bar to eject the molten alloy from the crucible nozzle.

The vertical rotating wheel is then operated to rotate at a selected tangential wheel velocity that is between about 3 m/s and about 30 m/s. While the wheel is rotating, a fluidic quenching/casting medium 42 is introduced into the space between the walls at the horizontal wheel face 38. Suitable fluidic media include liquids and gasses, e.g., water, whale oil, cottonseed oil, mineral oils, helium, chilled air, argon or other inert gas, or other selected liquid or gas. Additives such as polyalkylene glycol (PAG)-based synthetic products can be included. For many applications, water can be preferred as a quenching medium. As the drum wheel is rotated, the cooling medium circulates around the drum wheel.

The temperature of the quenching medium in the drum wheel can be actively controlled, e.g., to a temperature of between about 0° C. and about 100° C., for selected quenching media and selected processing applications. Such temperature control can be achieved by, e.g., a refrigeration or heating unit that cools or heats a selected quenching medium and feeds the temperature-controlled medium into the wheel. A selected quenching medium can be cooled or heated to achieve a desired melt casting operation, or the quenching medium can be selected for operation without active temperature control. For example, water as a quenching medium can be thermally controlled to a desired temperature that is above room temperature, or alternatively, unheated oil can be employed to achieve similar quenching results.

While a selected liquid quenching medium is continuously fed into the drum wheel at a selected tangential wheel velocity, the distance between the surface of the liquid and the lower end tip of the crucible nozzle is measured as that distance decreases due to the rising level of the liquid. When the distance between the nozzle tip and the liquid surface is between about 1 mm and about 50 mm, then the feed of liquid quenching media is terminated. The rotational speed of the wheel is then increased to a selected tangential wheel velocity between about 3 m/s and about 30 m/s. As explained in detail below, the wheel speed is preferably controlled based on a selected casting rate to achieve uniform casting structures, for example, to achieve a uniform wire diameter, by matching the wheel speed to the casting rate.

To begin melt spinning of the alloy material, the bulk solid alloy material in the crucible is melted, e.g., with induction coils around the crucible or with another suitable heating configuration. An inert gas, such as argon gas, is preferably continuously flowed through the crucible, out the nozzle, during this heating. A thermocouple or other suitable device can be disposed in the crucible with the alloy material to directly measure the temperature of the material during the heating process. Alternatively, an optical temperature reader or other device can be configured to sense and measure the alloy material temperature accurately from outside the crucible. No particular temperature measurement device is required. When the alloy material starts to melt and flow down through the nozzle to clog the nozzle, the flow of gas through the crucible is terminated and the crucible pressure is reduced to produce a vacuum, if desired. The temperature of the melting alloy material is then monitored. When the alloy material is fully melted and is at temperature that is between about 50° C. and about 500° C. above the alloy material melting temperature, defined herein as the liquidus temperature of the alloy composition in its phase diagram, the flow of gas is reintroduced to apply a pressure from the top of the crucible. The pressure flow is preferably sufficient to cause the melted alloy material to eject out of the crucible nozzle and into the quenching medium in the rotating drum.

As shown in FIG. 4, as the melted alloy material is ejected out of the nozzle 34, the alloy material takes on the cross-sectional geometry of the nozzle and forms a continuous structure such as a wire 10 that extends into and is collected by the rotating drum. Then as the structure enters the quenching medium, the alloy material solidifies into a continuous length of the cast geometry. To achieve continuous casting of significant lengths of alloy material, the wheel speed can be matched with the casting rate. The casting rate depends on the casting temperature, nozzle size, and pressure. For a given casting rate that results from these conditions, the wheel speed then is accordingly controlled. With wheel speed substantially matched to casting rate, the length of the cast geometry is limited only by the volume of alloy material that can be provided in the crucible. When all melted alloy material has been ejected from the crucible, the wheel rotation can be terminated, the quenching medium can be drained from the wheel, and the cast alloy structure can be collected from the wheel drum and wound or otherwise positioned. Other control parameters of the melt spinning process to be considered are prescribed in U.S. Pat. No. 10,167,540, issued Jan. 1, 2019, the entirety of which is hereby incorporated by reference.

The casting processes of FIGS. 3B and 3C are particularly advantageous in that very long lengths of cast structure, such as ferrous wire, can be uniformly produced. For example, in one embodiment herein, there is produced by a casting process a continuous ferrous wire that is longer than at least about one meter, and preferably that is longer than at least about 1.5 meters. The wire diameter along the length of the fiber is precisely controlled and as a result is highly uniform. The wire diameter uniformity is here specified for this embodiment as about ±5 microns along at least about a 1 meter-length of the wire. Tables I, II, and III specify three different embodiments provided herein of melt spinning processes that produce continuous, straight, substantially defect-free ferrous SMA wire or other structure.

TABLE I

Parameter
Range

Melt crucible nozzle diameter
30-1000
μm

Ejection pressure
1-10
bar

Tangential wheel velocity
3-30
m/s

Angle between jet and water
10-80°

Distance from nozzle tip to water
1-50
mm

Water temperature
0-100°
C.

Water layer thickness
>5
mm

TABLE II

Parameter
Range

Melt crucible nozzle diameter
50-300
μm

Ejection pressure
3-6
bar

Tangential wheel velocity
6-12
m/s

Angle between jet and water
20-60°

Distance from nozzle tip to water
3-30
mm

Water temperature
0-100°
C.

Water layer thickness
5-30
mm

TABLE III

Parameter
Range

Melt crucible nozzle diameter
100-200
μm

Ejection pressure
3-6
bar

Tangential wheel velocity
6-12
m/s

Angle between jet and water
25-50°

Distance from nozzle tip to water
3-30
mm

Water temperature
5-50°
C.

Water layer thickness
5-30
mm

Turning to the step of wire solutionization 115 in the method 100 of FIG. 3A, a length of ferrous wire resulting from the casting step is subject to a thermal solutionization step 115 for a duration that is sufficient to produce throughout the cross section and along the length of the wire a single phase of γ-fcc, and to cause the grains of the polycrystalline wire morphology to grow sufficiently large that the wire becomes a 1-D oligocrystalline structure. The solutionization process consists of subjecting the wire to an elevated temperature for a selected duration that produces a wire morphology which is a single phase γ matrix.

In various embodiments, the solutionization is conducted at a temperature between about 1100° C. and about 1500° C., for a duration of at least about one minute, and in various embodiments, the solutionization is conducted for a duration of about sixty minutes. In one embodiment, a ferrous alloy wire is subject to solutionization at a temperature of about 1300° C. for a duration of about sixty minutes or at least about sixty minutes. In a further embodiment, a ferrous alloy wire is subject to solutionization at a temperature of about 1300° C. for a duration of at least about one hour. In further embodiments, a ferrous alloy wire is subject to solutionization at a temperature of between about 1400° C. and about 1500° C. for a duration of at least one minute and less than about five minutes. In a further embodiment, a ferrous alloy wire is subjected to solutionization at a temperature of about 1200° C. for at least about 24 hours. In an example of a preferred embodiment, for a ferrous SMA wire of Fe-28Ni-17Co-11.5Al-2.5Ta-0.05B (atomic %), the wire is subjected to solutionization at a temperature of about 1300° C. for a duration of one hour.

Turning to the step of thermal ageing 120 in the method 100 of FIG. 3A, the length of oligocrystalline γ-fcc ferrous wire resulting from the solutionization step 115 is subject to a thermal ageing step 120 to form a dispersion of γ′-L1₂precipitates throughout the wire morphology. Specifically, the ageing process temperature and process time are selected to cause elements of the alloy to preferentially be in one of the two phases, either γ′ or γ. In various embodiments, the ageing is conducted at an elevated temperature, between about 400° C. and about 900° C., for an ageing duration of at least about one hour, and in various embodiments, is conducted for an ageing duration of at least about 24 hours, or at least 48 hours, and preferably for at least about 72 hours or a longer duration as-applicable to a given process and alloy composition.

The thermal ageing step 120 causes the precipitation of γ′-L1₂precipitates within the γ-fcc oligocrystalline matrix, resulting in a dispersion of such precipitates throughout the matrix. Uniform thermal ageing of the wire is preferred to maximize uniformity of the precipitate dispersion throughout the matrix. A population of precipitates that is at least about 10 volume % of the ferrous alloy can be preferred in various embodiments herein. The population of γ′-L1₂precipitates substantially enhances the strength of the oligocrystalline ferrous alloy wire, resulting in a wire tensile strength as high as 1.2 GPa or more, and substantially enhances the temperature range over which superelasticity is achievable for the oligocrystalline ferrous alloy wire. The dispersion of γ′-phase precipitates also enhances the repeatability of each shape-memory/superelastic cycle through which the oligocrystalline ferrous alloy wire is operated.

In the solutionization step 115 and ageing step 120 of the method sequence of FIG. 3A, it is preferred that the thermal processing be conducted in vacuum. Most ferrous alloy structures, such as a wire, will quickly oxidize otherwise. In preferred embodiments herein the heat treatments are conducted under a vacuum of about 8.5 mbar. This can be achieved in any convenient arrangement; for example, by encapsulating a wire in a quartz tube, or by other suitable arrangement. With such an arrangement, the wire can be heated under vacuum, and preferably heated by a thermal ramp, e.g., at a ramp rate of about 10° C./min. Each heat treatment is preferably ended by water quenching.

In embodiments herein, each and every step of the method sequence of FIG. 3A is conducted without solid state, i.e., mechanical, processing, or other processing that would increase the crystallographic texture of the ferrous alloy. In embodiments herein, any processing of the wire subsequent to the ageing step is also conducted without solid state processing that would increase the crystallographic texture of the ferrous alloy. In a preferred embodiment, no step of the wire production method includes solid state processing for increase in crystallographic texture. As explained above, it is discovered herein that such can be eliminated, thereby to increase efficiency and reduce cost of the production method, while at the same time achieving superior performance.

Thus, in one embodiment, grains of the γ-fcc matrix and γ′-L1₂precipitates within the γ-fcc oligocrystalline matrix are oriented at least partially randomly, with no little or no crystallographic texture imposed by any of the process steps. After thermal ageing of the ferrous SMA wire, the wire exhibits shape memory and superelasticity, and no further processing is required to achieve the thermoelastic properties herein. In other embodiments, the method of FIG. 3A is not limited to the process steps shown; other processing, such as diameter and surface condition tuning, can be conducted, both between the process steps and before/after the process sequence. But no solid state crystallographic texturing is to be conducted. The ferrous SMA wire exhibits superior shape memory and elasticity without crystallographic texturing.

The production method of FIG. 3A can be applied to any suitable structure, including tape, ribbon, or other cast structure. Other suitable processes can also be employed. It is to be recognized that for clarity, this descriptive example is specific to one structure, namely, a wire, but that the oligocrystalline morphology, ferrous composition, and phase composition can be applied to any suitable structure. Oligocrystalline ferrous pillars, tubes, sheets, cantilevers, walls, bridges, open cell foam, or other suitable geometric configuration can be employed with the matrix-precipitate morphology of the oligocrystalline ferrous material provided herein.

EXAMPLE

An ingot of Fe-28Ni-17Co-11.5Al-2.5Ta-0.05B (atomic %) alloy was fabricated by a vacuum arc re-melting (VAR) process using pure elements with >99.95 wt. % purity, except for FeB with a 99 wt. % purity. The ingot was re-melted five times and homogenized at 900° C. for 24 hours in an Ar atmosphere. Ferrous alloy wire was then cast by in rotating water melt-spinning (INROWASP), from which there were produced a length of wire that was sectioned into wire segments of varying lengths between about 10 cm and about 50 cm for purposes of experiment.

To this end, a single cuboidal piece of ingot material of about 3.5 g in weight was cut from the homogenized ingot and induction-melted in a quartz crucible under flowing Ar of 99.999% purity and then ejected by pressurized Ar into a rotating water wheel. Argon was ejected over the water surface for about 30 s and then shut off immediately before the melt was ejected through the gas blanket. The wire casting was conducted on an IN-Roquench machine with quartz crucibles (Phoenix Scientific Industries Ltd). Induction heating in the crucible was conducted with an Ambrell-EASYHEAT system operated at a frequency of 323 kHz. The current of the heating system was gradually increased at about 3·A·s⁻¹. The alloy material was observed to melt at a current of about 230 A, and the alloy was ejected when the current reached 320 A. The wire casting parameters are given in Table IV as follows:

TABLE IV

Parameter
Range

Melt crucible nozzle diameter
100
μm

Ejection pressure
5.5
bar

Wheel speed-530 mm-diameter
360
rpm

Angle between jet and water
35°

Distance from nozzle tip to water
20
mm

Water temperature
20°
C.

Water layer thickness
10
mm

The surface of the as-cast wire was observed to be reasonably smooth, without evidence of varicose, knot, or cracking defects. The wire cross-section was found to be elliptical, with an axial ratio of about 1.3±0.1. The as-cast microstructure of the wire was predominantly γ-phase, with some elemental macrosegregation in a cellular structure, and contained about 2.8% of β-NiAl phase precipitates. The average grain size was about 50 μm. The crystallographic texture of the as-cast wire was measured. FIG. 5 is an Inverse Pole Figure (IPF) showing the measured distribution of the orientation of grains in the wire. As-shown clearly by the IPF, no crystallographic texture is exhibited. The as-cast wire exhibited ductility of about 8% in tension, but did not exhibit superelasticity.

After casting, the ferrous wire was solutionized at a temperature of about 1300° C. for a duration of about 60 minutes. Thereafter the ferrous wire was aged at a temperature of about 600° C. for a duration of about 72 h. These solutionization and ageing conditions were determined to be the preferable conditions for producing a fine distribution of γ′ precipitates within the γ-phase matrix and for achieving an oligocrystalline morphology. For both of these heat treatments, the ferrous wire was encapsulated in a quartz ampoule under a vacuum of about 8.5 mbar. Three initial pump and purge cycles were first conducted with a gas flow Ar, and then the temperature ramped at a rate of about 10° C. min⁻¹. Once the process temperature was reached, the heating was held constant for the selected duration for each treatment. Then the wire was removed from the hot furnace and immediately quenched into room temperature water without breaking the quartz ampoules.

Separate lengths of the oligocrystalline ferrous wire were cut and arranged for microstructure characterization by hot-mounting in a conductive resin and mechanically polishing down to 1 μm diamond suspension followed by final polishing using 50 nm colloidal silica. The surface morphology and grain structure were characterized by scanning electron microscopy (SEM, JEOL 6610LV) and electron backscatter diffraction (EBSD, Zeiss Merlin). Needle-shaped samples for atom-probe tomography (APT) were prepared using a focused ion beam (FIB) lift-out technique in a FEI Helios NanoLab 450 F1 dual-beam FIB. APT experiments were performed using a local electrode atom probe (Cameca Instruments LEAP 4000X HR) in voltage mode at 50K. The pulse frequency and fraction were 200 kHz and 15%, respectively.

FIG. 6 is a 3-D reconstruction of the solutionized and aged oligocrystalline ferrous wire, produced from Atom Probe Tomography (APT). This reconstructed image confirms a distribution of γ′-L1₂precipitates each having an extent of between about 3 nm and about 10 nm, and a volume fraction of precipitates of about 30%.

After solutionizing and aging, there was found to be substantially no crystallographic texture present in the wire. FIG. 7 is an Inverse Pole Figure (IPF) showing the measured distribution of the orientation of grains in the finished wire. As-shown clearly by the IPF, no strong crystallographic texture is exhibited. FIG. 8 is an EBSD IFP map of the wire, confirming near-bamboo oligocrystalline grain structure, with one or two grains spanning the width of the wire. Most of the grain boundaries were identified as high-angle boundaries. Boundary precipitation of the β-NiAl phase is understood to embrittle the alloy and precludes superelasticity. Here, there is see no evidence of β-NiAl precipitation in the ferrous SMA wire, even at high-angle grain boundaries. Without being bound to theory, it is understood that this can be due to the boron content of the experimental wire, which was 0.07% measured by APT, which can be added in embodiments herein to suppress β-NiAl phase; and also is due to the general paucity of grain boundaries in the ferrous SMA wire.

Stress-strain curves were measured at constant temperature using a dynamic mechanical analyzer (DMA, TA Instruments 850) in load control mode at a constant rate of 5 N min⁻¹. Temperature control was provided by a gas flow system using liquid nitrogen. Shape-memory recovery strains were obtained in the DMA during a heating cycle to room temperature. Wire samples were gripped by aluminum foil tabs, which were bonded to the wires by a thin layer of epoxy (J-B Weld™). The gauge length of wire was between 6 and 8 mm. Besides cutting the ferrous SMA wire sections to length, no additional sample preparation was performed on the wire after the aging treatment. Strains were calculated from the crosshead displacement. This setup was confirmed to give accurate stiffness measurements on commercially available NiTi wires of similar diameter.

To generate the stress-strain curves, a length of the oligocrystalline ferrous wire was loaded to 700 MPa for 30 cycles, after which the maximum stress was increased on each subsequent cycle, until failure occurred, at a loading of 1.25 GPa, on the thirty eighth cycle. On the thirty seventh cycle, with the maximum stress of 1.2 GPa, a maximum superelastic strain of 7% was measured, corresponding to dissipated energy of 15.7 MJ m⁻³, during the full superelastic cycle. FIG. 9A is a plot of the measured engineering stress over the cycles.

To evaluate the repeatability of this performance, a length of the oligocrystalline ferrous wire was loaded for 20 stress-strain cycles each to 1 GPa, corresponding to about 4.5% strain. FIG. 9B is a plot of the measured stress over the series of cycles. As shown in the plot of FIG. 9B, on the first cycle, there were strain bursts, the total magnitude of which appears to correspond to the residual strain of about 1% after unloading, suggesting some plastic yielding in favorably oriented grains. After the exhaustion of these weak points in this training cycle, however, a closed superelastic loop was obtained on the next cycle. The subsequent stress-strain response was found to be quite stable, with a similar superelastic curve observed on the 20^thcycle with only a slight decrease, of between about 500 MPa and 400 MPa, in the stress to induce martensite, and a residual strain of 0.7%. This demonstrates that performing an initial training cycle at a high stress level can be an effective strategy to strain-harden the ferrous SMA wire material, enabling superelasticity up to elevated stresses of 1.2 GPa or higher.

The superelasticity performance shown in the plots of FIGS. 9A-9B confirm that the oligocrystalline ferrous wire exhibited superelastic behavior at room temperature. It is understood that this was enabled by precipitation hardening and solute partitioning during the thermal aging step, which led to a change in matrix composition of −11.3 at. % Ni, +2.7 at. % Co, and −4.3 at. % Al and an increase in the martensitic start temperature, M_s, of about 200° C. The aged wire demonstrated a maximum superelastic strain of 7% and tensile strength of about 1.2 GPa. This recovery strain level is among the highest reported superelastic strains in any untextured polycrystalline SMA, of any material composition, to date and is comparable to that of crystallographic textured NiTi wire.

Based on the measured data in the plots of FIGS. 9A-9B, it is shown that the stress hysteresis of the oligocrystalline ferrous wire provided herein, at between about 200 MPa and about 250 MPa, is considerably lower than the 500 MPa stress hysteresis that is conventionally expected of crystallographically-textured polycrystalline FeNiCoAlTaB. The stress hysteresis of the oligocrystalline ferrous wire provided herein is instead consistent with the stress hysteresis of between about 150 MPa and 200 MPa, that is conventionally observed in single crystal FeNiCoAlTaB sheets, and the stress hysteresis of between about 200 MPa and 300 MPa that is conventionally observed in crystallographically-textured polycrystalline NiTi wire. As such, the energy dissipated by the oligocrystalline ferrous wire provided herein, 15.7 MJ m⁻³, is less than the energy dissipated in crystallographically-textured polycrystalline FeNiCoAlTaB sheets, but similar to that of crystallographically-textured polycrystalline NiTi wires during a single superelastic cycle loading of 7% strain. This demonstrates that the reduced stress hysteresis of the oligocrystalline ferrous wire provided herein is directly beneficial for actuation applications and improves functional fatigue resistance, as is often the case in SMAs.

The superelastic curves of the plots of FIGS. 9A-9B do not exhibit a flat plateau corresponding to stress-induced transformation, but rather a steep incline, demonstrating significant transformation hardening at a high hardening rate of about 124 MPa %⁻¹. This is very significantly higher than the 11 MPa %⁻¹hardening rate that is expected for crystallographically-textured polycrystalline ferrous SMA and the 31 MPa %⁻¹hardening rate that is expected for single-crystal ferrous SMA. In addition, the plots of FIGS. 9A-9B demonstrate that the oligocrystalline ferrous wire provided herein exhibits a uniform transformation hardening rate throughout the entire transformation cycle and not an increasing hardening rate as can be expected for single-crystal material, which is attributed to martensite variant interactions accumulating with strain.

Without being bound to theory, it is understood herein that with the substantially random grain orientation distribution of the untextured oligocrystalline ferrous wire provided herein, some grains are relatively more favorably oriented for martensitic transformation and some grains are relatively less favorably oriented for transformation; as a result, different grains begin transforming to martensite at different macroscopic stresses. The polycrystalline microstructure of the oligocrystalline ferrous wire also exhibits more geometric constraints than that of single-crystal ferrous SMA, causing martensite variants to interact earlier in the transformation process and requiring additional stress to overcome the constraints and proceed with martensitic transformation. This effect is also likely enhanced by the small lattice symmetry change (cubic to tetragonal) of the γ-α′ transformation, which results in only 3 Lattice Correspondence Variants (LCVs) and 24 Habit Plane Variants (HPVs), compared to 12 LCVs and 192 HPVs for cubic-monoclinic NiTi. This limits the flexibility in transformation microstructures that can be formed to accommodate the imposed boundary conditions and may explain the significant transformation hardening of the oligocrystalline ferrous wire provided herein compared to other polycrystalline SMAs.

The superelasticity of the oligocrystalline ferrous wire provided herein was measured at temperatures between −40° C. and 30° C. FIG. 10A is a stress-strain plot for the oligocrystalline ferrous wire provided herein at each of the measured temperatures. FIG. 10B is a plot of the temperature dependence of martensite start (σM_s) and austenite finish (σ_Af) stresses, measured to be 5.4 MPa ° C.⁻¹. This is higher than the 2 MPa ° C.⁻¹-3 MPa ° C.⁻¹that is conventionally expected for [001] single-crystal FeNiCoAlTaB alloy. For a uniaxial strain, the Clausius-Clapeyron relationship states that

$\frac{d σ}{dT} = - \frac{Δ S}{ε},$

where ΔS is the transformation entropy and ε is the transformation strain. Thus, the smaller transformation strain in polycrystals can explain the higher

$\frac{d σ}{dT}$

compared to single crystals. However,

$\frac{d σ}{dT}$

in the oligocrystalline wire provided herein is still less than the 6-12 MPa ° C.⁻¹typically observed in NiTi, demonstrating that the ferrous SMA alloys wire provided herein is less temperature-sensitive than NiTi and can exhibit superelasticity over a wider temperature range than NiTi. Extrapolating σ_Afto zero stress gives an A_fvalue of −57° C.; thus, the material should be fully superelastic between −57° C. and at least 30° C., a temperature range of almost 100° C.

FIG. 11 is a plot of a stress-strain curve for the oligocrystalline ferrous wire provided herein loaded at a temperature of −100° C., exhibiting the double-yielding behavior typical of SMAs. The test temperature is slightly lower than the extrapolated M_sof −90° C., and thus the microstructure was here likely to consist of both austenite and martensite phases. There is virtually no initial elastic loading region exhibited before the first yield point, suggesting that the stress to induce additional martensite, and/or the martensite detwinning stress, is very low. The material exhibited linear elastic behavior during unloading, resulting in a permanent strain of 2.5%. As expected, the oligocrystalline ferrous wire is not superelastic at −100° C., but full recovery of this strain was achieved during heating back to 30° C. This demonstrates the classical one-way shape-memory property of aged FeNiCoAlTaB wires.

The Example described above experimentally confirms a 7% superelastic strain for the crystallographically untextured oligocrystalline ferrous SMA wire provided herein. This superelastic strain is much higher than that of conventional untextured polycrystalline ferrous SMA wire, which conventionally demonstrates early brittle fracture, and is competitive with that of crystallographically-textured polycrystalline ferrous SMA sheets and even single-crystal ferrous SMA sheets, both of which require considerably more intensive processing routes. Indeed, this recovery strain is among the highest reported to date for superelastic strains in any untextured polycrystalline SMA, of any composition, and is comparable to that of state-of-the-art, crystallographically-textured NiTi wire. This is particularly surprising because historically, it has been understood that in general in a SMA material, superelastic strain is significantly reduced in a random polycrystal grain orientation, compared with an organized polycrystal grain orientation; it was thought that the random grain orientation does not preferentially sample axes of large strain. The oligocrystalline ferrous SMA wire provided herein not only achieves a superelastic strain of 7% but also exhibits a tensile strength of about 1.2 GPa at room temperature. The oligocrystalline grain structure, dual phase morphology, and ferrous alloy composition operate synergistically to achieve this unexpected result even with crystallographic texturing of no more than that given for mrd=15. Along with excellent mechanical functionality, the oligocrystalline ferrous SMA wire provided herein is strongly ferromagnetic in the martensite phase, allowing for coupling of strain with magnetic field.

With this method, the oligocrystalline ferrous SMA wire provided herein can be produced by a rapid, continuous wire-casting process for producing kilometer lengths of wire. The combination of a high-throughput, low-cost processing methodology with substantially superior shape-memory properties enables widespread commercial application of ferrous shape-memory alloys. Additionally, SMA wire material that undergoes a small symmetry change, such as the cubic γ to tetragonal α′ transition in the oligocrystalline ferrous SMA wire provided herein, is conventionally understood to be an especially poor shape-memory material in a crystallographically-untextured form. The very unexpectedly and surprisingly high strain demonstrated in the untextured ferrous SMA wire provided herein bucks this historical convention and eliminates the cost and complexity of crystallographic texturing thought to be required to achieve high superelastic strain. Further, the cyclic stability of the oligocrystalline ferrous SMA wire provided herein is considerably better than that of crystallographically-textured polycrystalline FeNiCoAlTiB wire and compares favorably to single-crystal FeNiCoTaAl wire and NiTi wire at similar strain amplitudes. This enables the oligocrystalline ferrous SMA wire to be employed with confidence of long time-to-failure at small strain amplitudes, a condition for which NiTi wire has conventionally been preferred.

It is recognized that those skilled in the art may make various modifications and additions to the embodiments described above without departing from the spirit and scope of the present contribution to the art. Accordingly, it is to be understood that the protection sought to be afforded hereby should be deemed to extend to the subject matter claims and all equivalents thereof fairly within the scope of the invention.

Superelastic and Shape-Memory Fe-Based Wire and Direct Wire Casting Thereof

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