Metallic alloy compositions and methods of forming the same

Abstract

A metallic alloy system having good strength and ductility is provided. For example, a Fe-XAl-YTi alloy system for use in low and high temperature applications can include one or more elements to enhance properties and / or performance of the metal alloy at high temperatures, while also having excellent yield strength and ultimate tensile strength properties. Example embodiments of the alloy system can include Fe-29.98Ni-14Co-XAl-2Ta-YTi-5V-0.02B, where X+Y=12, and Fe-31Ni-15Co-7Al-2Ti-5V-2Ta-0.02B. The atomic percentages of the elements in the compositions of these alloy systems can be adjusted to change stability of the phases thereof.

Description

Attorney Docket No.: MIT 26210 PCT | 88212-432499METALLIC ALLOY COMPOSITIONS AND METHODS OF FORMING THE SAMECROSS REFERENCE TO RELATED APPLICATION(S)

[0001] The present disclosure claims priority7to and the benefit of U.S. Provisional Application No. 63 / 742,246. entitled ‘“Metallic Alloy Compositions." filed on January 6. 2025, the content of which is incorporated by reference herein in its entirety.FIELD

[0002] The present disclosure relates to metallic alloys, and more particularly relates to Fe-XAl-YTi alloy systems exhibiting dual-phase multi -precipitate microstructures that achieve combinations of high yield strength and high ductility7.BACKGROUND

[0003] Metallic alloys serve as foundational materials across a wide range of industrial applications, including automotive, aerospace, energy generation, and defense sectors. The performance of structural components in these applications depends on the mechanical properties of the alloys from which they are fabricated, with strength and ductility being among the most significant characteristics. Conventional alloy design strategies have sought to balance these properties, as increases in strength are often accompanied by reductions in ductility, and vice versa.

[0004] Traditional approaches to achieving high strength in steel alloys have relied on the introduction of martensite through rapid cooling from elevated temperatures. Martensite contributes to strength through its high dislocation density, boundary density, and trapped interstitial carbon content. However, martensitic microstructures exhibit limited capacity for plastic deformation and can serve as sources for micro-crack nucleation during forming operations or at higher deformation levels. Efforts to enhance ductility in steel alloys have typically involved the incorporation of metastable austenite, which can undergo mechanically induced martensitic transformation to produce transformation-induced plasticity effects. While this approach can delay necking and provide toughening benefits, the eventual transformation to fresh martensite limits the extent of mechanical benefits and can lead to micro-crack nucleation.

[0005] Bey ond the mechanical limitations associated with martensite and metastable austenite, additional challenges arise in processing and utilization of alloys containing theseAttorney Docket No.: MIT 26210 PCT | 88212-432499microstructural constituents. High cooling rates are ty pically seen in martensite formation, while achieving optimal austenite stability to maximize transformation-induced plasticity¬ effects requires careful optimization of composition and microstructuralfeatures. Establishing microstructure-property relationships in alloys that utilize phase metastability can be complex, as mechanical behavior can vary depending on processing conditions. This complexity can result in performance of dedicated analyses for different processing recipes and alloy grades.

[0006] Precipitation-strengthening approaches have been explored as alternatives to martensite-based and metastability-based design strategies. These approaches have included single matrix systems with single precipitate types, such as face-centered cubic matrices with LI2 precipitates or body-centered cubic matrices with B2 precipitates. However, the intrinsic brittleness of ordered precipitate phases can exacerbate strength-ductility tradeoffs. Microstructural heterogeneity and the high Peierls stress of ordered intermetallic compounds can lead to strain accumulation at ordered-disordered interfaces, potentially resulting in early micro-crack formation and plastic instability.

[0007] Accordingly, there is a need for alloy compositions and microstructural architectures that can achieve combinations of high strength and high ductility without relying on martensite or metastable austenite phases, which can provide sustained strain hardening behavior through the coordinated contributions of multiple microstructural constituents.SUMMARY

[0008] The presently disclosed embodiments generally relate to metallic alloy compositions and methods of forming thereof, particularly Fe-Ni-Co-Al-Ti-V-Ta-B alloy systems that exhibit a dual-phase multi-precipitate microstructure. Such alloy systems achieve combinations of high strength and high ductility without retying on martensite or metastable austenite phases. The disclosed alloy compositions include face-centered cubic (fee) grains containing LI2 precipitates and B2 grains containing body-centered cubic (bcc) precipitates, wi th the hierarchical microstructure enabling sustained strain hardening behavior through coordinated contributions of multiple microstructural constituents. Some nonlimiting examples of the compositions can include Fe-29.98Ni-14Co-XAl-2Ta-YTi-5V-0.02B, where X+Y=12, and Fe-31Ni-15Co-7Al-2Ti-5V-2Ta-0.02B. The atomic percentagesAttorney Docket No.: MIT 26210 PCT | 88212-432499of the elements in the compositions of these alloy systems can be adjusted to change stability of the phases thereof.

[0009] In accordance with some embodiments, the ability to tune the relative fractions of fee and bcc phases through compositional adjustments, particularly by varying aluminum and titanium content, can present advantages in tailoring mechanical properties for specific applications. The thermomechanical processing routes disclosed herein, including annealing, cold rolling, recry stallization, and aging treatments, enable the development of multi-phase microstructures that exhibit multiple stages of strain hardening during deformation. These alloy systems can be suitable for use in applications requiring high strength and ductility at temperatures ranging from room temperature to intermediate elevated temperatures.

[0010] In an aspect, embodiments relate to a composition comprising an alloy. The alloy has iron (Fe) approximately in a range of about 35 at.% to about 40 at.%, cobalt (Co) approximately in a range of about 12 at.% to about 17 at.%, nickel (Ni) approximately in a range of about 28 at.% to about 32 at.%. aluminum (Al) approximately in a range of about 5 at.% to about 14 at.%, tantalum (Ta) approximately in a range of about 1 at.% to about 2.5 at.%, titanium (Ti) approximately in a range of about 0 at.% to about 7 at.%, vanadium (V) approximately in a range of about 3 at.% to about 8 at.%, and boron (B) approximately in a range of about 0.01 at.% to about 0.3 at.%. A total composition of Al and Ti is approximately in a range of about 5 at.% to about 14 at.%.

[0011] One or more of the following features can be included. The alloy composition can be stable at a temperature approximately in a range of about 550 degrees Celsius to about 800 degrees Celsius. The Ni can be approximately in a range of about 29.98 at.% to about 31 at.%. The Co can be approximately in a range of about 14 at.% to about 15 at.%. The Al can be approximately in a range of about 7 at.% to about 10 at.%. The Ta can be at about 2 at.%. The Ti can be at about 2 at.%. The V can be at about 5 at.%. The B can be at about 0.02 at.%. The B can be included as a master alloy of Fe-B to prevent evaporation during casting.

[0012] In another aspect, embodiments relate to a method of forming an alloy. The method comprises combining an iron-aluminum-titanium (Fe-Al-Ti) alloy system with a cobalt (Co) approximately in a range of about 12 at.% to about 17 at.%, a nickel (Ni) approximately in a range of about 28 at.% to about 32 at.%. a tantalum (Ta) approximately in a range of about 1 at.% to about 2.5 at.%, a vanadium (V) approximately in a range of about 3 at.% to about 8 at.%, and a boron (B) approximately in a range of about 0.01 at.% to about 0.3 at.%, to formAttorney Docket No.: MIT 26210 PCT | 88212-432499an alloy system. A total composition of Al and Ti is approximately in a range of about 5 at.% to about 14 at.%.

[0013] One or more of the following features can be included. The method can further comprise adjusting an amount of one or more of the Co, Ni, Ta, V, or B to change phase stability of the alloy system. The method can further comprise adjusting an amount of one or more of Al or Ti in the alloy system. Adjusting can further comprise increasing the amount of Al or decreasing the amount of Ti. The Ni can be approximately in a range of about 29.98 at.% to about 31 at.%. The Co can be approximately in a range of about 14 at.% to about 15 at.%. The B can be added as a master alloy of Fe-B to prevent evaporation during casting. The Al can be approximately in a range of about 7 at.% to about 10 at.%. The Ta can be at about 2 at.%. The Ti can be at about 2 at.%. The V can be at about 5 at.%. The B can be at about 0.02 at.%. The total composition of Al and Ti can be about 12 at.%. The alloy system can be stable at a temperature approximately in a range of about 550 degrees Celsius to about 800 degrees Celsius.

[0014] In another aspect, embodiments relate to a composition comprising an alloy that includes one or more of a face-centered cubic phase that contains Lb precipitates distributed therein, or a B2 phase that contains body-centered cubic precipitates distributed therein.

[0015] One or more of the following features can be included. The alloy can comprise a dual-phase matrix architecture composed of the face-centered cubic phase and the B2 phase. The face-centered cubic phase can further contain B2 precipitates distributed therein. The alloy can be free from metastable austenite phases and martensite phases. The Lh precipitates can be coherent with the face-centered cubic phase. A lattice misfit between the LI2 precipitates and the face-centered cubic phase can be approximately -0.84%. The bodycentered cubic precipitates can be coherent with the B2 phase. A lattice misfit between the body-centered cubic precipitates and the B2 phase can be approximately -0.26%. The bodycentered cubic precipitates within the B2 phase can have a cuboidal morphology. The alloy can exhibit at least three distinct strain hardening stages during tensile deformation. The alloy can be formed from a master alloy of iron and boron combined with additional alloying elements. The additional alloying elements can include nickel, cobalt, aluminum, titanium, tantalum, and vanadium.

[0016] In another aspect, embodiments relate to a method of forming an alloy. The method comprises combining iron, nickel, cobalt, aluminum, and vanadium to form an alloy system.Attorney Docket No.: MIT 26210 PCT | 88212-432499The method further comprises subj ecting the alloy system to a recry stallization treatment. The method further comprises subjecting the alloy system to an aging treatment to precipitate Lh precipitates within face-centered cubic grains and body-centered cubic precipitates within B2 grains, thereby forming a dual-phase matrix architecture.

[0017] One or more of the following features can be included. Combining iron, nickel, cobalt, aluminum, and vanadium can further comprise combining yvith titanium, tantalum, and boron. The boron can be included as a master alloy of iron-boron to prevent evaporation during casting. The recrystallization treatment can be performed at a temperature in a range of about 1150 degrees Celsius to about 1190 degrees Celsius. The recrystallization treatment can be performed for a duration in a range of about 15 seconds to about 20 minutes. The aging treatment can be performed at a temperature of approximately 650 degrees Celsius. The aging treatment can be performed for a duration in a range of about 10 hours to about 30 hours. The method can further comprise cold rolling the alloy system yvith approximately 60 percent reduction in thickness prior to the recrystallization treatment. The method can further comprise homogenizing the alloy system at approximately 1150 degrees Celsius for approximately 10 hours prior to cold rolling. The alloy can further comprise iron in a range of about 35 at.% to about 40 at.%. nickel in a range of about 28 at.% to about 32 at.%, cobalt in a range of about 12 at.% to about 17 at.%, aluminum in a range of about 5 at.% to about 14 at.%, and vanadium in a range of about 3 at.% to about 8 at.%. The alloy can further comprise titanium in a range of about 0 at.% to about 7 at.% and tantalum in a range of about 1 at.% to about 2.5 at.%.BRIEF DESCRIPTION OF DRAWINGS

[0018] This disclosure will be more fully understood from the following detailed description taken in conjunction with the accompanying drawings, in which:

[0019] FIG. 1A is an isopleth pseudo-binary’ phase diagram showing temperature versus mole fraction of aluminum for FeNi3oCoi4AlxTii2-xV5Ta2 alloys calculated using thermodynamic modeling software (ThermoCalc 2025a, TCFE13 database);

[0020] FIG. IB is a phase fraction versus temperature plot for an Fe-6Al-6Ti composition;

[0021] FIG. 1C is a phase fraction versus temperature plot for an Fe-8Al-4Ti composition;

[0022] FIG. ID is a phase fraction versus temperature plot for an Fe-10Al-2Ticomposition;Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0023] FIG. IE is a phase fraction versus temperature plot for an Fe-12A1 composition;

[0024] FIG. 2A is an Ashby plot comparing ultimate tensile strength versus total elongation for various steel types and high-entropy alloy systems;

[0025] FIG. 2B is an Ashby plot comparing ultimate tensile strength versus total elongation for single precipitate alloys and multi-precipitate alloy compositions;

[0026] FIG. 3A is an engineering stress versus engineering strain plot comparing alloy systems having single-matrix, single-precipitate and a designed dual-phase, multi -precipitate phase constitution;

[0027] FIG. 3B is an engineering stress versus engineering strain plot for multiple alloy compositions within an Fe-XAl-YTi system;

[0028] FIG. 4 is a table summarizing mechanical properties for various Fe-XAl-YTi alloy compositions subjected to different heat treatments;

[0029] FIG. 5 is a schematic diagram representing a heat-treatment schedule and processing route for alloy design;

[0030] FIG. 6 is a graph showing strain hardening rate as a function of true strain for four different alloy compositions;

[0031] FIG. 7A is a strain hardening rate versus true strain plot for an Fe-6Al-6Ti alloy showing multistage strain hardening behavior;

[0032] FIG. 7B is a strain hardening rate versus true strain plot for an Fe-8Al-4Ti alloy showing multistage strain hardening behavior;

[0033] FIG. 7C is a strain hardening rate versus true strain plot for an Fe-10Al-2Ti alloy showing multistage strain hardening behavior;

[0034] FIG. 7D is a strain hardening rate versus true strain plot for an Fe-12A1 alloy showing multistage strain hardening behavior;

[0035] FIG. 8 is a strain hardening rate versus true strain plot for an Fe-10Al-2Ti alloy with annotations indicating contributions of different phases to each strain hardening stage;

[0036] FIG. 9A is an electron backscatter diffraction phase map for an Fe-6Al-6Ti alloy composition;Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0037] FIG. 9B is an electron backscatter diffraction phase map for an Fe-8Al-4Ti alloy composition;

[0038] FIG. 9C is an electron backscatter diffraction phase map for an Fe-10Al-2Ti alloy composition;

[0039] FIG. 9D is an electron backscatter diffraction phase map for an Fe-12A1 alloy composition;

[0040] FIG. 10A is a schematic diagram illustrating wavy slip deformation mode in alloys during uniaxial tensile loading;

[0041] FIG. 10B is a schematic diagram illustrating planar slip deformation mode in alloys during uniaxial tensile loading;

[0042] FIG. 10C is a schematic diagram illustrating high density dislocation walls deformation mode in alloys during uniaxial tensile loading;

[0043] FIG. 10D is a schematic diagram illustrating microbands deformation mode in alloys during uniaxial tensile loading;

[0044] FIG. 11 A is an electron channeling contrast imaging micrograph of an Fe-10Al-2Ti alloy at a plastic strain of 0.05 showing planar slip features;

[0045] FIG. 1 IB is an electron channeling contrast imaging micrograph of the Fe-10Al-2Ti alloy of FIG. 11A showing high density dislocation walls;

[0046] FIG. 11 C is an inverse pole figure map showing deformation mechanisms at a plastic strain of 0.05 for the Fe-10Al-2Ti alloy of FIG. 11A;

[0047] FIG. 12A is an electron channeling contrast imaging micrograph of an Fe-10Al-2Ti alloy at a plastic strain of 0.15 showing high density dislocation walls;

[0048] FIG. 12B is an electron channeling contrast imaging micrograph of the Fe-10Al-2Ti alloy of FIG. 12A showing microbands;

[0049] FIG. 12C is an inverse pole figure map showing deformation mechanisms at a plastic strain of 0.15 for the Fe-10Al-2Ti alloy of FIG. 12A;

[0050] FIG. 13A is an electron channeling contrast imaging micrograph of an Fe-10Al-2Ti alloy at a plastic strain of 0.26 showing microbands;Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0051] FIG. 13B is an inverse pole figure map showing deformation mechanisms at a plastic strain of 0.26 for the Fe-10Al-2Ti alloy of FIG. 13 A;

[0052] FIG. 14A is an electron channeling contrast imaging micrograph of an Fe-10Al-2Ti alloy at a plastic strain of 0.395 showing microbands;

[0053] FIG. 14B is an inverse pole figure map showing deformation mechanisms at a plastic strain of 0.395 for the Fe-10Al-2Ti alloy of FIG. 14A;

[0054] FIG. 15 is a backscattered scanning electron microscope micrograph of an aged Fe-10Al-2Ti alloy showing hierarchical microstructure with identified regions;

[0055] FIG. 16A is the micrograph of FIG. 15 having a face-centered cubic (fee) phase marked;

[0056] FIG. 16B is an electron diffraction pattern along a

[0011] zone axis showing Lb superlattice spots alongside fundamental fee spots for the alloy of FIG. 1 A;

[0057] FIG. 16C is an electron diffraction pattern along a

[0112] zone axis showing Lb superlattice spots alongside fundamental fee spots for the alloy of FIG. 16A;

[0058] FIG. 16D is a dark-field transmission electron microscopy image showing distribution of Lb precipitates for the alloy of FIG. 16 A;

[0059] FIG. 17A is the micrograph of FIG. 15 having a B2 phase marked;

[0060] FIG. 17B is a selected area electron diffraction pattern along a

[0001] zone axis showing B2 superlattice spots alongside fundamental bcc spots for the alloy of FIG. 17A;

[0061] FIG. 17C is a higher magnification transmission electron microscopy image showing cuboidal bcc precipitates within a B2 matrix for the alloy of FIG. 17 A;

[0062] FIG. 17D is a dark-field transmission electron microscopy image showing distribution of bcc precipitates within a B2 matrix for the alloy of FIG. 17A;

[0063] FIG. 18 is an X-ray diffraction pattern for an Fe-10Al-2Ti alloy composition showing FCC-Lb and BCC-B2 phase peaks;

[0064] FIG. 19A is a one-dimensional compositional profile along a Z-axis for a B2 grain with bcc precipitates from atom probe tomography analysis shown in FIG. 19B;

[0065] FIG. 19B is a three-dimensional reconstruction of a B2 grain region with bcc precipitates from atom probe tomography analysis;Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0066] FIG. 19C is a three-dimensional reconstruction of an fee grain region with LI 2 precipitates from atom probe tomography analysis;

[0067] FIG. 19D is a one-dimensional compositional profile of FIG. 19C;

[0068] FIG. 20 is an X-ray diffraction characterization graph showing intensity versus diffraction angle for Fe-6Al-6Ti, Fe-8Al-4Ti, and Fe-12A1 alloys;

[0069] FIG. 21A is a weak-beam dark-field transmission electron microscopy image showing planar slip with precipitate shearing in fee grains of an Fe-10Al-2Ti alloy;

[0070] FIG. 21B is a bright-field transmission electron microscopy image showing dislocation features in the Fe-10Al-2Ti alloy of FIG. 21 A;

[0071] FIG. 21C is a dark-field transmission electron microscopy image showing dislocation features in the Fe-10Al-2Ti alloy of FIG. 21 A;

[0072] FIG. 22A is a weak-beam dark-field transmission electron microscopy image show ing dislocation activity within a B2 phase of an Fe-10Al-2Ti alloy;

[0073] FIG. 22B is a higher magnification weak-beam dark-field transmission electron microscopy image showing bcc precipitate shearing in the Fe-10Al-2Ti alloy of FIG. 22A;

[0074] FIG. 23 A is the backscattered scanning electron microscope micrograph of FIG. 15 showing B2 grains inside the fee grains;

[0075] FIG. 23B is a selected area electron diffraction pattern from a dark phase of an fee grain showing bcc and B2 reflections for the alloy of FIG. 23A;

[0076] FIG. 23C is a transmission electron microscopy dark-field image showing bcc precipitates within a B2 matrix for the alloy of FIG. 23 A;

[0077] FIG. 23D is a transmission electron microscopy dark-field image showing a B2 matrix with bcc precipitates for the alloy of FIG. 23 A;

[0078] FIG. 24A is an electron channeling contrast imaging micrograph of an Fe-6Al-6Ti alloy deformed to 13% plastic strain showing wavy slip deformation;

[0079] FIG. 24B is a higher magnification electron channeling contrast imaging micrograph of the Fe-6Al-6Ti alloy of FIG. 24A showing wavy slip features;

[0080] FIG. 25A is an electron channeling contrast imaging micrograph of an Fe-8Al-4Ti alloy deformed to 5% plastic strain showing wavy slip deformation;Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0081] FIG. 25B is an electron channeling contrast imaging micrograph of the Fe-8Al-4Ti alloy of FIG. 25 A showing planar slip deformation;

[0082] FIG. 25C is an electron channeling contrast imaging micrograph of the Fe-8Al-4Ti alloy of FIG. 25A deformed to 17% plastic strain showing wavy slip;

[0083] FIG. 25D is an electron channeling contrast imaging micrograph of the Fe-8Al-4Ti alloy of FIG. 25A deformed to 17% plastic strain showing high density dislocation walls;

[0084] FIG. 26 A is an electron channeling contrast imaging micrograph of an Fe-12A1 alloy deformed to 5% plastic strain showing wavy slip deformation;

[0085] FIG. 26B is an electron channeling contrast imaging micrograph of the Fe-12A1 alloy of FIG. 26A showing planar slip deformation;

[0086] FIG. 26C is an electron channeling contrast imaging micrograph of the Fe-12A1 alloy of FIG. 26A showing high density dislocation walls formation;

[0087] FIG. 26D is an electron channeling contrast imaging micrograph of the Fe-12A1 alloy of FIG. 26A deformed to 22% plastic strain showing high density dislocation walls;

[0088] FIG. 26E is an electron channeling contrast imaging micrograph of the Fe-12A1 alloy of FIG. 26A deformed to 22% plastic strain showing microbands formation;

[0089] FIG. 27A is a backscattered electron scanning electron microscope micrograph of a recrystallized Fe-10Al-2Ti alloy showing dual-phase grain structure;

[0090] FIG. 27B is an engineering stress versus engineering strain curve for the recrystallized Fe-10Al-2Ti alloy of FIG. 27 A;

[0091] FIG. 27C is a strain hardening rate versus true strain plot for the recrystallized Fe-10AI-2Ti alloy of FIG. 27 A;

[0092] FIG. 27D is a transmission electron microscope bright-field image showing microbands in the deformed recrystallized Fe-10Al-2Ti alloy of FIG. 27A;

[0093] FIG. 27E is a transmission electron microscope image showing high density dislocation walls in the deformed recrystallized Fe-10Al-2Ti alloy of FIG. 27 A; and

[0094] FIG. 27F is an inverse pole figure map showing deformation mechanisms at 35% plastic strain for the recrystallized Fe-10Al-2Ti alloy of FIG. 27A.Attorney Docket No.: MIT 26210 PCT | 88212-432499DETAILED DESCRIPTION

[0095] Certain embodiments will now be described to provide an overall understanding of the principles of the structure, function, manufacture, and use of the systems, devices, related components (e.g, alloy compositions, microstructural phases, precipitates, matrix grains, and thermomechanical processing equipment), and methods disclosed herein. One or more examples of these embodiments are illustrated in the accompanying drawings. Those skilled in the art will understand that the devices and methods specifically described herein and illustrated in the accompanying drawings are non-limiting exemplary embodiments and that the scope of the present disclosure is defined solely by the claims. The features illustrated or described in connection with one exemplary’ embodiment may be combined with the features of other embodiments. Such modifications and variations are intended to be included within the scope of the present disclosure. Further, to the extent features, phases, stages, steps, or the like are described as being "first." "second," "third," etc., and / or "lower," "upper," "middle," etc., such numerical and / or location ordering / identifi cation is generally arbitrary, and thus such numbering can be interchangeable unless indicated or otherwise understood by those skilled in the art to not be interchangeable.

[0096] The figures provided herein are not necessarily to scale, although a person skilled in the art will recognize instances where the figures are to scale and / or what a ty pical size is when the drawings are not to scale. Further, to the extent that linear or circular dimensions or shapes are used or described herein, such dimensions are not intended to limit the types of shapes or sizes of such devices, components, etc. A person skilled in the art will recognize that an equivalent to such linear and / or circular dimensions or shapes can be easily- determined for any geometric shape (e.g, references to widths and diameters being easily adaptable for circular and linear dimensions, respectively, by a person skilled in the art). While in some embodiments movement of one component is described with respect to another, a person skilled in the art will recognize that other movements are possible. Further, to the extent arrows are used to describe a direction a component can expand or move, these arrows are illustrative and in no way limit the direction the respective component can expand or move. A person skilled in the art will recognize other ways and directions for creating the desired tension or movement.

[0097] Still further, in the present disclosure, like-numbered components of various embodiments generally have similar features when those components are of a similar nature and / or serve a similar purpose, unless otherwise noted or otherwise understood by a personAttorney Docket No.: MIT 26210 PCT | 88212-432499skilled in the art. To the extent the present disclosure includes prototy pes, mock-ups, bench models, or the like, a person skilled in the art will recognize how to rely upon the present disclosure to integrate the techniques, systems, devices, and methods into a product, such as structural components for automotive, aerospace, energy generation, and defense applications. A number of terms may be used throughout the disclosure interchangeably but will be understood by a person skilled in the art. By way of non-limiting example, the terms "alloy," "alloy system," and "alloy composition" may be used interchangeably with one another, the terms "face-centered cubic," "fee," and "y phase" may be used interchangeably with one another, the terms “y'” and “L12”may be used interchangeably with one another, and the terms "body-centered cubic," "bcc," and "a phase" may be used interchangeably with one another. Moreover, it will be appreciated that although features may be discussed with respect to one embodiment within the present disclosure, these features can be applied to every7embodiment of the present disclosure where such feature would be supported.

[0098] To the extent terms like "approximately," "about," and "substantially" are used herein, a person skilled in the art will appreciate the scope those words convey in the context of their usage. In the context of alloy compositions and microstructural characterization, obtaining a certain degree of compositional precision, phase fraction measurement, precipitate size distribution, and / or lattice parameter determination, among other measurements and the like, may be difficult, and thus use of terms like "approximately," "about," and "substantially" is intended to address this difficulty'. A person skilled in the art will understand what constitutes how close a particular dimension or composition should be to still fall within the spint of the quantification and description provided for herein. Even in instances where such terminology is not used, and a dimension or composition just includes the number or term (e.g., "coherent" is used instead of "substantially coherent"), a person skilled in the art will appreciate that, unless explicitly indicated otherwise, terms like "approximately," "about," and "substantially" are applicable to those dimensions and terms as well. The foregoing notwithstanding, a person skilled in the art will appreciate that terms like "approximately," "about," and "substantially" at least encompass compositions, dimensions, temperatures, and quantities that are ±10%, 10°, etc. of the provided amount, or encompass dimensions that are ±5%, 5°, etc. of the provided amount, unless indicated otherwise or otherwise known to those skilled in the art. The present disclosure appreciates that a person skilled in the art, in view of the present disclosure, understands suitable compositions and processing parameters for various features of the disclosed alloy systems, microstructures,Attorney Docket No.: MIT 26210 PCT | 88212-432499and related components of any of the same, and thus to the extent a particular composition or processing parameter is described, unless it is explicitly indicated that such composition or processing parameter is required, a person skilled in the art will appreciate other compositions or processing parameters that are possible without impacting the overall alloy system or microstructure.

[0099] The present disclosure relates to alloy compositions and alloy systems that achieve high strength and high ductility without relying on martensite or metastable austenite phases. Conventional approaches to improving the strength-ductility trade-off in steels and high-entropy alloys have relied on introducing metastability in the austenitic matrix or incorporating martensitic phases. However, such design strategies can incur penalties in the form of premature damage nucleation and early failure during deformation, and such strategies can suffer from limited strength retention at elevated temperatures. The alloy compositions and alloy systems of the present disclosure address these limitations by employing a dual-phase multi-precipitate microstructure architecture that avoids deformation-induced martensitic transformation.

[0100] A composition of the present disclosure can include an alloy having a dual -phase matrix architecture. The dual-phase matrix architecture can be composed of a face-centered cubic (fee) phase and a B2 phase, where the fee phase and the B2 phase serve as matrices, or grains, for various nanoscale precipitates. It will be appreciated that the term "B2 phase" refers to an ordered body-centered cubic structure of the CsCl-type, and the term "fee phase" refers to a disordered face-centered cubic structure. The fee phase of the alloy can contain nanoscale LH precipitates and body-centered cubic (bcc) precipitates distributed at the nanoscale, and B2 precipitates at the microscale. It will be appreciated that the term "LI 2" refers to an ordered face-centered cubic structure of the CmAu-type. which is characteristic of y' precipitates in nickel-based and cobalt-based superalloys. The B2 phase of the alloy can contain bcc nanoscale precipitates distributed therein. The hierarchical distribution of multiple precipitate types within the dual-phase matrix architecture can provide multiple stages of strain hardening during deformation, thereby enabling the alloy to achieve both high strength and high ductility.

[0101] The alloy of the present disclosure can be free from metastable austenite phases and martensite phases. By avoiding metastable austenite and martensite, the alloy can avoid deformation-induced martensitic transformation during mechanical loading. The absence of deformation-induced martensitic transformation can reduce the likelihood of prematureAttorney Docket No.: MIT 26210 PCT | 88212-432499damage nucleation that is associated with phase transformation during deformation. The multi-precipitate microstructure embedded within the dual-phase matrix architecture can provide strain hardening through dislocation-precipitate interactions rather than through phase transformation mechanisms. The alloy compositions of the present disclosure can thus achieve strength-ductility combinations that exceed those of conventional steels containing martensite or metastable austenite, as well as those of single-precipitate alloy systems.

[0102] The alloy compositions of the present disclosure can be designed based on valence electron concentration (VEC) considerations for phase stability. The VEC can be calculated according to the equation:

[0103] where Ci is the atomic fraction of element i and (VEC)i is the valence electron concentration of element i. Within this framework, VEC > 8.0 can stabilize the fee phase, whereas 6.7 < VEC < 8.0 can favor the coexistence of fee and bcc phases. The alloy design strategy' of the present disclosure can employ VEC-based predictions to tune the relative fractions of the fee and bcc phases by varying the concentrations of alloying elements.

[0104] Thermodynamic phase diagrams and phase fraction plots illustrate the evolution of phases in Fe-Ni-Co-based alloys. In particular, FIGS. 1 A- IE illustrate the evolution of phases of FeNisoCouAlxTi -xVsTaz alloys as a function of aluminum content and temperature. FIG. 1 A shows an isopleth phase diagram with temperature on the vertical axis and mole fraction of aluminum on the horizontal axis. The isopleth phase diagram illustrates various phase regions including fee, fee + Lb, B2 + fee, bcc + fee, bcc + fee + Lb, and bcc + B2 + fee + Lb phase fields. As shown in FIG. 1A, increasing aluminum content can promote the formation of the B2 phase and the bcc phase at the expense of the fee phase, consistent with VEC-based predictions.

[0105] With continued reference to FIGS. 1A-1E, FIGS. IB through IE show phase fraction versus temperature plots for alloy compositions having different aluminum and titanium contents for a subset of the region of FIG. 1A for alloys having Fe-XALYTi, where X+Y=12, with 12 indicating an atomic percentage (at.%) can show optimal strain hardening rates. While strain hardening regimes can be manipulated by changing the constitution of the phases and their fractions, Fe-XAl-YTi, where X+Y=12 presents optimal mechanical properties, such as engineering stress vs strain results, as discussed in greater detail below. Observation of the isopleth binary phase diagram in FIG. 1 A shows a narrow region betweenAttorney Docket No.: MIT 26210 PCT | 88212-4324990.04 to 0.12 mole fraction of Al where the dual-phase, multi-precipitate BCC + B2 + FCC + L12 can be stable between a temperature of 600 degrees Celsius to 800 degrees Celsius. Fe-XAl-YTi alloys having compositions: i) Fe-6Al-6Ti, ii) Fe-8Al-4Ti, iii) Fe-10Al-2Ti, and iv) Fe-12A1 were selected and aged at 650 degrees Celsius such that the thermodynamic calculations can be confirmed for temperature boundaries accuracy in this range. FIG. IB shows a phase fraction plot for the Fe-6Al-6Ti composition, FIG. 1C shows a phase fraction plot for the Fe-8Al-4Ti composition, FIG. ID shows a phase fraction plot for the Fe-lOAl-2Ti composition, and FIG. IE shows a phase fraction plot for the Fe-12A1 composition. As illustrated in FIGS. IB through IE, the volume fraction of B2 + bcc phases increases with increasing aluminum content in the alloy. The Fe-6Al-6Ti composition shown in FIG. IB exhibits a predominantly fee microstructure, whereas the Fe-12A1 composition shown in FIG. IE exhibits a larger B2 phase fraction.

[0106] The alloy compositions of the present disclosure can include nickel (Ni) and cobalt (Co) as fee stabilizers. Ni can be present in a range of about 28 at.% to about 32 at.%, and Co can be present in a range of about 12 at.% to about 17 at.%. Ni and Co can stabilize the fcc-y phase of the alloy, enabling the precipitation of LG-ordered precipitates within the stabilized fee phase. Because Ni preferentially partitions to y'-LE precipitates of the NijAl ty pe, with a partitioning tendency approximately twice that of Co in y'-lNi.Co.FefiAl precipitates, the Ni content can be maintained at least twice that of Co. The Ni concentration can be limited to preserve an Fe-rich composition and to avoid suppressing bcc phase formation, as increasing Ni content raises the VEC and stabilizes the fee phase.

[0107] The alloy compositions of the present disclosure can include vanadium (V) and aluminum (Al) as bcc stabilizers and LE stabilizers. V can be present in a range of about 3 at.% to about 8 at.%, and Al can be present in a range of about 5 at.% to about 14 at.%. V can stabilize the bcc phase in the Fe system and can stabilize LE precipitates through selective partitioning to the y' precipitates. Al can promote bcc phase stability while also contributing to LE phase stability through preferential partitioning. The substitution of V and Al in place of heavier bcc stabilizers such as tungsten (W), molybdenum (Mo), and chromium (Cr) can reduce the mass density of the alloy. Mo and Cr can destabilize the LE phase and can promote the formation of deleterious topologically close-packed (TCP) phases, such as p and o phases.

[0108] The alloy compositions of the present disclosure can include tantalum (Ta) and titanium (Ti) as LE stabilizers. Ta can be present in a range of about 1 at.% to about 2.5Attorney Docket No.: MIT 26210 PCT | 88212-432499at.%, and Ti can be present in a range of about 0 at.% to about 7 at.%. Ta and Ti can stabilize the Lk-ordered precipitates within the fee matrix, similar to y' precipitates in Ni-based superalloys. Ta can be a more potent y' former than niobium (Nb), and the Ta concentration can be fixed at about 2 at.% because the solubility7of Ta in Fe is limited. Additions of Ta beyond 2 at.% in CoNi-based alloys can promote TCP phase formation. The total composition of Al and Ti can be in a range of about 5 at.% to about 14 at.%.

[0109] The alloy compositions of the present disclosure can include iron (Fe) as the balance element. Fe can be present in a range of about 35 at.% to about 40 at.%. The alloy compositions can also include boron (B) in a range of about 0.01 at.% to about 0.3 at.%, or in a range of about 0.01 at.% to about 0.03 at%. B additions can enhance grain boundary7cohesion through B-metal hybridization at grain boundaries. B can form covalent bonds with surrounding metal atoms, enhancing cohesion without depleting metal-metal bonding electrons. The B content can be limited to avoid the formation of Ta-rich borides at grain boundaries.

[0110] A method of forming an alloy of the present disclosure can include combining an iron-aluminum-titanium (Fe-Al-Ti) alloy system with Co in a range of about 12 at.% to about 17 at.%, Ni in a range of about 28 at.% to about 32 at.%, Ta in a range of about 1 at.% to about 2.5 at.%. V in a range of about 3 at.% to about 8 at.%, and B in a range of about 0.01 at.% to about 0.3 at.% to form an alloy system. The total composition of Al and Ti in the alloy system can be in a range of about 5 at.% to about 14 at.%. The method can further include adjusting an amount of one or more of Co, Ni, Ta, V, or B to change phase stability7of the alloy system. The method can further include adjusting an amount of one or more of Al or Ti in the alloy system. Adjusting can include increasing the amount of Al or decreasing the amount of Ti, which can increase the volume fraction of B2 + bcc phases at the expense of the fee phase, as illustrated in FIGS. IB through IE.

[0111] Referring to FIGS. 2A-2B, Ashby plots illustrate ultimate tensile strength versus total elongation for various alloy systems, demonstrating the mechanical property7advantages of the dual-phase multi -precipitate architecture of the present disclosure. In particular, FIG.2A shows data for different steel types including TRIP / metastable high entropy7alloys.TRIPLEX steels, dual-phase steels, martensitic steels, and TRIP maraging steels. A data point corresponding to the fee + bcc + LL + B2 composition of the present disclosure is positioned at approximately 1600 MPa ultimate tensile strength and approximately 45 percent total elongation, demonstrating performance beyond the conventional alloy groupings shownAttorney Docket No.: MIT 26210 PCT | 88212-432499as shaded elliptical regions in FIG. 2A. The alloy compositions of the present disclosure can thus achieve strength-ductility combinations that exceed those of conventional steels containing martensite or metastable austenite.

[0112] With continued reference to FIGS. 2A-2B. FIG. 2B compares tensile strength in single precipitate alloys with the alloy compositions of the present disclosure. For example, the shaded regions in FIG. 2B represent single precipitate fee + LG alloys and fcc / bcc + B2 alloys, and compares them to the fee + bcc + Lb + B2 compositions of the present disclosure. Moreover, the data points in FIG. 2B show the progression from fee + LG compositions to fee + LG + B2 compositions, and to the fee + bcc + LG + B2 compositions of the present disclosure. As shown, the fee + bcc + LG + B2 compositions achieve approximately 50 percent total elongation at approximately 1900 MPa ultimate tensile strength. Analysis of the dashed diagonal lines in both FIGS. 2A and 2B represent constant strength-ductility product contours, and the alloy compositions of the present disclosure can achieve strength-ductility products that exceed those of single-precipitate alloy systems.

[0113] The Fe-10Al-2Ti composition of the present disclosure, which is relied upon throughout the present disclosure when describing the fee + bcc + LG + B2 compositions due to its desirable characteristics, many of which will be discussed in greater detail below, can achieve an ultimate tensile strength of approximately 1650 ± 60 MPa and an elongation-to-failure of approximately 38 ± 3%. The Fe-10Al-2Ti composition can achieve a yield strength of approximately 1050 ± 50 MPa. The Fe-10Al-2Ti composition can exhibit a UTS-to-YS ratio of approximately 1.57. The UTS-to-YS ratio of approximately 1.57 indicates that the alloy can undergo substantial strain hardening between the yield point and the ultimate tensile strength, which can be attributed to the multiple stages of strain hardening provided by the hierarchical multi -precipitate microstructure. The combination of high ultimate tensile strength, high yield strength, and high elongation-to-failure can enable the alloy compositions of the present disclosure to be used in applications requiring both high strength and high ductility at temperatures approximately in a range of about 550 degrees Celsius to about 800 degrees Celsius. It will be appreciated that the ultimate tensile strength and the elongation-to-failure can have an inverse relationship such that an increase in ultimate tensile strength can result in a decrease in elongation-to-failure and vice versa. In some embodiments, for example, the ultimate tensile strength of the composition of the present disclosure can be approximately 2100 MPa, while achieving 15% elongation-to-failure, with said composition being contemplated within the scope of the present disclosure. Further details about theAttorney Docket No.: MIT 26210 PCT | 88212-432499possible properties of the compositions as related to ultimate tensile strength and elongation-to-failure are discussed below with respect to FIG. 4.

[0114] Referring to FIGS. 3A-3B, engineering stress-strain curves illustrate the mechanical behavior of various Fe-Al-Ti compositions and demonstrate the advantages of the dual-phase multi -precipitate architecture over single-precipitate systems. FIG. 3A shows engineering stress in MPa on the vertical axis and engineering strain in percent on the horizontal axis for three alloy systems having different phase constitutions. A first curve in FIG. 3A labeled bcc + B2 corresponds to an Fe-Ni-Al-Mo-Nb-C-B alloy and reaches approximately 2000 MPa ultimate tensile strength with limited elongation of approximately 10 percent. A second curve in FIG. 3A labeled fee + Lb corresponds to an Fe-Ni-ALTi-Zr-C-B alloy and reaches approximately 1400 MPa with elongation of approximately 40 percent. A third curve in FIG.3 A corresponds to the Fe-10Al-2Ti composition of the present disclosure having fee + bcc + LL + B2 phase constitution and achieves approximately 1650 MPa ultimate tensile strength with approximately 38 percent elongation. As illustrated in FIG. 3A, the Fe-10Al-2Ti composition achieves a superior strength-ductility' balance compared to single-precipitate systems, combining the high strength characteristics of bcc + B2 alloys with the high ductility characteristics of fee + Lb alloys.

[0115] The desirable characteristics of the Fe-10Al-2Ti composition as compared to the remaining Fe-XALYTi composition of the present disclosure can be discussed with continued reference to FIGS. 3A-3B. For example, FIG. 3B shows engineering stress in MPa on the vertical axis and engineering strain in percent on the horizontal axis for multiple alloy compositions within the Fe-XALYTi system of the present disclosure. The Fe-10AL2Ti composition achieves the highest ultimate tensile strength of approximately’ 1650 MPa with elongation of approximately 38 percent. The Fe-8Al-4Ti composition reaches approximately 1500 MPa with approximately 22 percent elongation. The Fe-12A1 composition achieves approximately 1360 MPa with approximately 23 percent elongation. The Fe-6Al-6Ti composition reaches approximately 1200 MPa with approximately 13 percent elongation. An inset graph in FIG. 3B compares the Fe-10Al-2Ti alloy with an Fe-Co-Ni-ALTi-B alloy having fee + Lb microstructure, demonstrating the advantage of the multi-precipitate dualphase microstructure over single precipitate systems in Fe-Ni-Co containing alloy systems.

[0116] Referring to FIG. 4, a table summarizes alloy system mechanical properties for various Fe-based alloy compositions subjected to different heat treatments. FIG. 4 presents composition, heat treatment conditions, yield strength at 25 degrees Celsius in megapascals,Attorney Docket No.: MIT 26210 PCT | 88212-432499ultimate tensile strength at 25 degrees Celsius in megapascals, and total elongation as a percentage for each alloy composition. The mechanical properties data presented in FIG. 4 demonstrate the range of strength-ductility' combinations achievable through the alloy compositions and thermomechanical processing methods of the present disclosure.

[0117] It will be appreciated that while the present disclosure discusses alloy systems having Fe-XAl-YTi, where X+Y=12, other compositions can fall within the scope of the present disclosure. For example, as show n in FIG. 4, the alloy compositions of the present disclosure can include a composition variant designated Fe-7Al-2Ti having the formula Fe-31Ni-15Co-7Al-2Ta-2Ti-5V-0.02B. The Fe-7Al-2Ti composition can exhibit a single-phase fee microstructure w ithout the B2 phase. The absence of the B2 phase in the Fe-7Al-2Ti composition can be attributed to the reduced aluminum content compared to the Fe-10Al-2Ti composition, which can maintain the VEC above the threshold for B2 phase formation. The Fe-7Al-2Ti composition can achieve a yield strength in a range of approximately 1070 MPa to approximately' 1500 MPa, an ultimate tensile strength in a range of approximately 1680 MPa to approximately 2010 MPa, and an elongation in a range of approximately 23 percent to approximately 53 percent, depending on heat treatment conditions. The variation in mechanical properties of the Fe-7Al-2Ti composition with heat treatment conditions demonstrates that the alloy compositions of the present disclosure can be tuned to achieve different strength-ductility' combinations through thermomechanical processing.

[0118] The alloy' compositions of the present disclosure can be stable at a temperature approximately in a range of about 550 degrees Celsius to about 800 degrees Celsius. The temperature stability of the alloy compositions can be attributed to the LE precipitates, which can have a solvus temperature of approximately 900 degrees Celsius. The Lb precipitates can remain stable during service at temperatures below the solvus temperature, thereby maintaining the strengthening effect of the precipitates during elevated temperature applications. A method of use of the alloy compositions of the present disclosure can include employing the alloy system in applications where the alloy system is stable at a temperature approximately in a range of about 550 degrees Celsius to about 800 degrees Celsius. Such applications can include advanced ultra-supercritical pow er plants, nuclear applications, and automotive, military7, and aerospace applications requiring high strength and high ductility7at elevated temperatures.

[0119] As shown in FIG. 4, the Fe-12A1 composition having the formula Fe-30Ni-14Co-12Al-5V-2Ta-0.02B can be heat treated at 1150 degrees Celsius for 20 minutes followed byAttorney Docket No.: MIT 26210 PCT | 88212-432499aging at 650 degrees Celsius for 10 hours. The Fe-12A1 composition subjected to this heat treatment can exhibit a yield strength of 910 ± 20 megapascals, an ultimate tensile strength of 1360 ± 20 megapascals, and a total elongation of 23 ± 2 percent. The reduced ultimate tensile strength of the Fe-12A1 composition compared to the Fe-10Al-2Ti composition can be attributed to the higher B2 phase fraction, which can result in a decrease in the strengthductility combination when the B2 phase fraction exceeds an optimal range. The Ni content of approximately 30 at.% in the Fe-12A1 composition falls within the range of about 29.98 at.% to about 31 at.%, and the Co content of 14 at.% falls within the range of about 14 at.% to about 15 at.%.

[0120] With continued reference to FIG. 4, the Fe-10Al-2Ti composition having the formula Fe-30Ni-14Co-10Al-5V-2Ti-2Ta-0.02B can be subjected to various heat treatment conditions to achieve different mechanical property combinations. When solutionized at 1150 degrees Celsius for 20 hours, the Fe-10Al-2Ti composition can exhibit a yield strength of 580 ± 10 megapascals, an ultimate tensile strength of 915 ± 10 megapascals, and a total elongation of 31 ± 3 percent. When heat treated at 1150 degrees Celsius for 20 minutes without aging, the Fe-10Al-2Ti composition can exhibit a yield strength of 570 ± 10 megapascals, an ultimate tensile strength of 970 ± 20 megapascals, and a total elongation of 36 ± 1 percent.

[0121] As further shown in FIG. 4, when the Fe-10Al-2Ti composition is heat treated at 1150 degrees Celsius for 20 minutes followed by aging at 650 degrees Celsius for 10 hours, the yield strength can increase to 1050 ± 50 megapascals, the ultimate tensile strength can reach 1630 ± 60 megapascals, and the total elongation can be 38 ± 3 percent. The increase in yield strength and ultimate tensile strength upon aging can be attributed to the precipitation of LL precipitates within the fee grains and bcc precipitates within the B2 grains during the aging heat treatment. When heat treated at 1190 degrees Celsius for 20 minutes followed by aging at 650 degrees Celsius for 10 hours, the Fe-10Al-2Ti composition can exhibit a yield strength of 800 ± 30 megapascals, an ultimate tensile strength of 1380 ± 30 megapascals, and a total elongation of 32 ± 2 percent. The Ni content of approximately 30 at.% in the Fe-10Al-2Ti composition falls within the range of about 29.98 at.% to about 31 at.%, and the Co content of 14 at.% falls within the range of about 14 at.% to about 15 at.%.

[0122] With continued reference to FIG. 4, the Fe-8Al-4Ti composition having the formula Fe-30Ni-14Co-8Al-5V-4Ti-2Ta-0.02B can be heat treated at 1150 degrees Celsius for 20 minutes followed by aging at 650 degrees Celsius for 10 hours. The Fe-8Al-4Ti compositionAttorney Docket No.: MIT 26210 PCT | 88212-432499subjected to this heat treatment can exhibit a yield strength of 900 ± 40 megapascals, an ultimate tensile strength of 1400 ± 50 megapascals, and a total elongation of 22 ± 4 percent. The increased elongation of the Fe-8Al-4Ti composition compared to the Fe-6A1-6Ti composition can be attributed to the presence of the B2 phase and bcc precipitates that provide additional strain hardening mechanisms.

[0123] The Fe-6Al-6Ti composition having the formula Fe-30Ni-14Co-6Al-2Ta-5V-6Ti-0.02B can be subjected to the same heat treatment of 1150 degrees Celsius for 20 minutes followed by aging at 650 degrees Celsius for 10 hours. The Fe-6Al-6Ti composition can exhibit a yield strength of 905 ± 30 megapascals, an ultimate tensile strength of 1300 ± 50 megapascals, and a total elongation of 13 ± 2 percent. The limited elongation of the Fe-6A1-6Ti composition can be attributed to the absence of the B2 phase and bcc precipitates that contribute to additional strain hardening stages in the Fe-10Al-2Ti composition.

[0124] The Fe-7Al-2Ti composition having the formula Fe-31Ni-15Co-7Al-5V-2Ti-2Ta-0.02B can be subjected to various heat treatment conditions as shown in FIG. 4. When heat treated at 1150 degrees Celsius for 20 minutes without aging, the Fe-7Al-2Ti composition can exhibit a yield strength of 520 ± 20 megapascals, an ultimate tensile strength of 880 ± 30 megapascals, and a total elongation of 45 ± 2 percent. When heat treated at 1150 degrees Celsius for 5 minutes followed by aging at 650 degrees Celsius for 20 hours, the Fe-7Al-2Ti composition can exhibit a yield strength of 1115 ± 10 megapascals, an ultimate tensile strength of 1635 ± 20 megapascals, and a total elongation of 39 ± 3 percent. When heat treated at 1150 degrees Celsius for 1 minute followed by aging at 650 degrees Celsius for 20 hours, the Fe-7Al-2Ti composition can exhibit a yield strength of 1070 ± 20 megapascals, an ultimate tensile strength of 1680 ± 30 megapascals, and a total elongation of 53 ± 1 percent. When heat treated at 1150 degrees Celsius for 45 seconds followed by aging at 650 degrees Celsius for 20 hours, the Fe-7Al-2Ti composition can exhibit a yield strength of 1110 ± 30 megapascals, an ultimate tensile strength of 1780 ± 10 megapascals, and a total elongation of 49 ± 3 percent. When heat treated at 1150 degrees Celsius for 15 seconds followed by aging at 650 degrees Celsius for 20 hours, the Fe-7Al-2Ti composition can exhibit a yield strength of 1500 ± 50 megapascals, an ultimate tensile strength of 2010 ± 50 megapascals, and a total elongation of 23 ± 3 percent. The Ni content of approximately 31 at.% in the Fe-7Al-2Ti composition falls within the range of about 29.98 at.% to about 31 at.%, and the Co content of 15 at.% falls within the range of about 14 at.% to about 15 at.%.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0125] The mechanical properties data presented in FIG. 4 demonstrate that a method of forming an alloy of the present disclosure can include combining an iron-aluminum-titanium (Fe-Al-Ti) alloy system with Ni in a range of about 29.98 at.% to about 31 at.% and Co in a range of about 14 at.% to about 15 at.%. The method can further include subjecting the alloy to heat treatment conditions including solutionizing at temperatures in a range of about 1150 degrees Celsius to about 1190 degrees Celsius for durations ranging from about 15 seconds to about 20 hours, followed by aging at 650 degrees Celsius for durations ranging from about 10 hours to about 20 hours. The alloy compositions of the present disclosure can maintain mechanical stability at 650 degrees Celsius, as evidenced by temporal evolution of Vickers hardness during aging at 650 degrees Celsius. The stability of Vickers hardness during aging can indicate that the precipitate microstructure remains stable during extended exposure to elevated temperatures, thereby enabling the alloy compositions to be used in applications requiring sustained mechanical performance at temperatures in a range of about 550 degrees Celsius to about 800 degrees Celsius. For example, in some embodiments, at 550 degrees Celsius, the material can maintain a yield strength of about 770 MPa. with an ultimate tensile strength of about 950 MPa, while at 650 degrees Celsius, the material can maintain a yield strength of about 770 MPa, with an ultimate tensile strength of about 710 MPa.

[0126] Referring to FIG. 5, a schematic diagram illustrates a heat-treatment schedule and processing route utilized in the design of the alloy compositions of the present disclosure. The thermomechanical processing route can include multiple sequential steps to develop the hierarchical multi-precipitate microstructure that provides the strength-ductility combination of the alloy compositions. The processing sequence shown in FIG. 5 begins with an annealing step, followed by water quenching, cold rolling, recrystallization, water quenching, aging, and a final water quenching step.

[0127] The alloy compositions of the present disclosure can be processed by vacuum arcmelting pure metals. Following arc-melting, the alloy can be suction cast into a copper mold to form rectangular ingots. The rectangular ingots can have dimensions of approximately 20 x 8 x 70 mm3. The suction casting into a copper mold can provide rapid solidification of the alloy, which can refine the as-cast microstructure and reduce segregation of alloying elements.

[0128] Boron can be added to the alloy using a master alloy of iron and boron (Fe-B) to prevent boron evaporation during casting. The use of the Fe-B master alloy can ensure that the boron content is retained in the alloy during the high-temperature arc-meltingAttorney Docket No.: MIT 26210 PCT | 88212-432499process. The boron content in the as-cast alloy can range from approximately 0.07 at.% to approximately 0.26 at.% depending on the casting method used. The variation in boron content with casting method can be attributed to differences in cooling rate and exposure time at elevated temperatures during solidification.

[0129] With continued reference to FIG. 5, the processing route can include homogenizing the as-cast alloy at approximately 1150°C for 10 hours. The homogenization treatment, also referred to as the annealing step in FIG. 5, can reduce compositional segregation that develops during solidification. Following the homogenization treatment, the alloy can be water quenched to retain the high-temperature microstructure.

[0130] The processing route can include cold rolling the homogenized alloy with approximately 60% reduction in thickness. As shown in FIG. 5, the cold rolling can be performed in steps from about 8 mm to about 3.2 mm thickness. The cold rolling can introduce stored strain energy into the alloy, which can drive recrystallization during subsequent heat treatment. The 60% reduction in thickness can provide sufficient stored strain energy to achieve complete recrystallization of the alloy during the recrystallization treatment.

[0131] Following cold rolling, the alloy can be subjected to a recrystallization treatment at temperatures between about 1150°C and about 1190°C. The recrystallization temperature can be varied between 1150°C, 1175°C, and 1190°C to tune the microstructure and properties of the alloy. The recrystallization treatment can be performed for durations ranging from 15 seconds to 20 minutes. As illustrated in FIG. 5, the recrystallization treatment can be followed by water quenching to retain the recrystallized microstructure. The variation in recrystallization temperature and duration can affect the grain size and phase distribution of the alloy, which can influence the mechanical properties.

[0132] The processing route can include an aging treatment at approximately 650°C following the recrystallization treatment. The aging treatment can be performed at about 650°C for durations ranging from 10 hours to 20 hours. As shown in FIG. 5, the aging treatment can be performed for about 20 to about 30 hours in some embodiments. The aging treatment can precipitate the Lb precipitates within the fee grains and the bcc precipitates within the B2 grains, thereby developing the hierarchical multi-precipitate microstructure. Following the aging treatment, the alloy can be water quenched to retain the aged microstructure. The combination of recrystallization and aging treatments can enableAttorney Docket No.: MIT 26210 PCT | 88212-432499the tuning of microstructural phase constituents, which can govern the different regimes of strain-hardening behavior during deformation and can result in the high strength and ductility properties exhibited by the alloy compositions of the present disclosure.ro 1331 Referring to FIG. 6, a graph illustrates strain hardening rate as a function of true strain for four different alloy compositions of the present disclosure. The vertical axis of FIG. 6 represents strain hardening rate measured in units of 103MPa, ranging from 0 to 15, while the horizontal axis of FIG. 6 represents true strain expressed as a percentage, ranging from 0 to 35 percent. Four data series are plotted in FIG. 6 corresponding to the four different alloy compositions of the present disclosure, with Fe-6Al-6Ti shown with triangle markers, Fe-8Al-4Ti shown with circle markers, Fe-10Al-2Ti shown with square markers, and Fe-12A1 shown with diamond markers.

[0134] As shown in FIG. 6, the Fe-6Al-6Ti alloy exhibits a strain hardening rate that decreases from approximately 6 103MPa to approximately 3 * 103MPa over a true strain range of approximately 2 percent to approximately 13 percent before the curve terminates. The Fe-8Al-4Ti alloy shows a similar initial strain hardening rate that decreases from approximately 6 x 103MPa to values below 3 * 103MPa over a true strain range extending to approximately 25 percent. The Fe-12A1 alloy exhibits strain hardening behavior with the rate decreasing from approximately 12 x 103MPa to approximately 1 x 103MPa over a true strain range of approximately 2 percent to approximately 17 percent.

[0135] The alloy compositions of the present disclosure can exhibit four distinct strain hardening rate (SHR) stages during tensile deformation. The four distinct SHR stages can enable sustained work hardening and delayed plastic instability during tensile loading. The sustained w ork hardening provided by the four SHR stages can delay the onset of necking instability, thereby enabling the alloy to achieve high uniform elongation before failure. The delayed plastic instability can be attributed to the sequential activation of different strain hardening mechanisms associated with the Lh precipitates, the B2 grains, and the B2 + bcc phases within the fee grains at different strain levels.

[0136] With continued reference to FIG. 6, the Fe-10Al-2Ti composition demonstrates strain hardening behavior extending to approximately 35 percent true strain, with the strain hardening rate decreasing from approximately 6 x 103MPa to approximately 2 x 103MPa over this range. The Fe-10Al-2Ti composition maintains strain hardening capability over the largest range of true strain among the four compositions tested. The extended strainAttorney Docket No.: MIT 26210 PCT | 88212-432499hardening range of the Fe-10Al-2Ti composition can be attributed to the hierarchical multiprecipitate microstructure that provides multiple mechanisms for dislocation storage and strain hardening throughout the deformation process.[01371 The strain hardening regimes of the alloy compositions of the present disclosure can be manipulated by changing the phase constitution and phase fractions. As illustrated in FIG.6, the Fe-6Al-6Ti composition having a predominantly fee + LL microstructure exhibits a limited strain hardening range compared to the Fe-10Al-2Ti composition having the fee + bcc + Lh + B2 microstructure. The Fe-8Al-4Ti and Fe-12A1 compositions exhibit intermediate strain hardening ranges that correspond to the intermediate B2 phase fractions in these compositions. The ability to manipulate strain hardening regimes by changing phase constitution and fractions can enable the alloy compositions of the present disclosure to be tuned for different mechanical property requirements.

[0138] The variation in strain hardening behavior among the four alloy compositions shown in FIG. 6 can be correlated with the phase fractions discussed with reference to FIGS. IB through IE. The Fe-6Al-6Ti composition having the lowest B2 phase fraction exhibits the shortest strain hardening range, whereas the Fe-10Al-2Ti composition having an intermediate B2 phase fraction exhibits the longest strain hardening range. The Fe-12A1 composition having the highest B2 phase fraction exhibits a reduced strain hardening range compared to the Fe-10Al-2Ti composition, indicating that an optimal B2 phase fraction exists for maximizing the strain hardening range. The optimal B2 phase fraction can provide a balance between the strain hardening contributions from the LE precipitates in the fee grains and the bcc precipitates in the B2 grains.

[0139] The SHR stages during tensile deformation are discussed in greater detail with respect to FIGS. 7A-7D, and the discussion of each stage of strain hardening rate is included with respect to the precipitates discussed in FIG. 8. For example, referring to FIGS. 7A-7D, strain hardening rate versus true strain plots illustrate multistage strain hardening behavior in the alloy compositions of the present disclosure (shown in FIG. 6) during tensile loading at room temperature. FIG. 7 A shows the strain hardening rate as a function of true strain for the Fe-6Al-6Ti alloy, displaying two distinct stages of strain hardening labeled as Stage 1 and Stage 2. Stage 2 in FIG. 7A spans approximately 10 percent of the true strain range. The Fe-6Al-6Ti alloy having a predominantly fee + LE microstructure exhibits a limited number of strain hardening stages compared to the dual-phase multi-precipitate compositions.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0140] With continued reference to FIGS. 7A-7D, FIG. 7B shows the strain hardening rate versus true strain for the Fe-8Al-4Ti alloy, exhibiting three stages of strain hardening. Stage 2 in FIG. 7B covers approximately 10 percent of the true strain range, and Stage 3 covers approximately 10 percent of the true strain range. FIG. 7D shows the strain hardening rate versus true strain for the Fe-12A1 alloy, also exhibiting three stages of strainhardening. Stage 2 in FIG. 7D covers approximately 3 percent of the true strain range, and Stage 3 covers approximately 12 percent of the true strain range. The Fe-8Al-4Ti and Fe-12A1 alloys having dual -phase fee + B2 microstructures exhibit an additional strain hardening stage compared to the Fe-6Al-6Ti alloy, which can be attributed to the presence of the B2 grains and bcc precipitates in the dual-phase compositions.

[0141] FIG. 7C shows the strain hardening rate versus true strain for the Fe-10Al-2Ti alloy, displaying four distinct stages of strain hardening. Stage 2 in FIG. 7C spans approximately 6.5 percent of the true strain range, Stage 3 spans approximately 11 percent of the true strain range, and Stage 4 spans approximately 12.5 percent of the true strain range. The Fe-lOAl-2Ti alloy exhibits the largest number of strain hardening stages among the four compositions, which can be attributed to the hierarchical multi -precipitate microstructure containing LL precipitates in the fee grains, bcc precipitates in the B2 grains, and B2 + bcc phases within the fee grains.

[0142] Referring to FIG. 8, a strain hardening rate versus true strain plot for the Fe-lOAl-2Ti alloy includes annotations indicating the contributions of different phases to each strain hardening stage. As shown in FIG. 8, the LI 2 Phase region corresponds to early strain values, the B2 Grains region corresponds to intermediate strain values, and the B2 Phase of FCC Grains region corresponds to higher strain values extending to approximately 35 percent true strain. The annotations in FIG. 8 illustrate that each phase and precipitate type in the hierarchical microstructure can contribute to a distinct strain hardening stage during tensile deformation.

[0143] The Lb precipitates within the fee grains can govern Stage 2 of the strain hardening behavior. During Stage 2, dislocations can interact with the LG precipitates through shearing mechanisms, which can provide strain hardening through the creation of anti-phase boundaries (APBs) within the precipitates. The Ti / Al ratio in the alloy compositions can control the anti-phase boundary energy of the LI2 precipitates. The anti-phase boundary energy can affect the shearability of the Lk precipitates, which can influence the deformation mechanisms operating during Stage 2. An increase in Ti content can reduce the shearabilityAttorney Docket No.: MIT 26210 PCT | 88212-432499of the LI2 precipitates due to the higher planar fault energy of the precipitates, which can inhibit planar slip and can lead to the formation of dislocation walls and microbands that can degrade the mechanical properties.[01441 The B2 grains containing bcc precipitates can contribute to Stage 3 of the strain hardening behavior. During Stage 3, the B2 grains can accommodate plastic strain through dislocation activity' within the B2 matrix and through shearing of the bcc precipitates. The co-deformability between the fee grains and the B2 grains can enable continued strain hardening during Stage 3 without premature damage nucleation at the interphase boundaries. The transition from Stage 2 to Stage 3 can occur when the strain hardening contribution from the LL precipitates in the fee grains begins to saturate and the B2 grains begin to contribute to the overall strain hardening response.

[0145] With continued reference to FIG. 8, the microscale B2s within the fee grains can control Stage 4 of the strain hardening behavior through microband formation. The B2s within the fee grains can cause dislocation pileups along the primary slip system during deformation. The dislocation pileups can lead to stress localization at the B2s, which can promote cross-slip of dislocations to secondary and tertiary slip systems. The cross-slip events can facilitate microbanding, where dislocations from multiple slip systems interact to form dense dislocation bands aligned along crystallographic planes. The microbands can provide long-range back stress that can sustain the strain hardening rate during Stage 4, thereby delaying plastic instability' and enabling the alloy to achieve high uniform elongation. Visualization of each of the slips, microbands, and dislocations will be included with respect to the figures below.

[0146] FIGS. 9A-9D illustrate electron backscatter diffraction (EBSD) phase maps 100 of the microstructural evolution of the alloy compositions of the present disclosure as aluminum content increases. FIGS. 9A-9D display face-centered cubic (fee) regions 102 and bodycentered cubic (bcc) regions 104 for each of the four compositions of the present embodiments. The EBSD phase maps 100 demonstrate how the phase constitution of the alloy compositions changes with variations in aluminum and titanium content, consistent with the thermodynamic predictions discussed with reference to FIGS. 1 A-1E.

[0147] FIG. 9A corresponds to the Fe-6Al-6Ti alloy composition and shows a predominantly fee microstructure with large grains 102 and minimal bcc phase content. The Fe-6Al-6Ti composition can exhibit a single-phase fcc-based microstructure, which can beAttorney Docket No.: MIT 26210 PCT | 88212-432499atributed to the VEC of approximately 8.0 that stabilizes the fee phase. The grain size distribution of the Fe-6Al-6Ti alloy can exhibit a mode at approximately 200 pm to approximately 350 pm, representing the coarsest grains among the alloy compositions of the present disclosure. The large grain size of the Fe-6Al-6Ti composition can be attributed to the absence of the B2 phase, which can otherwise provide a grain-pinning effect during recrystallization.

[0148] With continued reference to FIGS. 9A-9D, FIG. 9B corresponds to the Fe-8Al-4Ti alloy composition and displays a finer grain structure with small dispersed bcc phase regions 104 within the fee matrix 102. The introduction of the B2 phase in the Fe-8Al-4Ti composition can result from the increased aluminum content, which can lower the VEC below' 8.0 and can stabilize the fee + bcc phase field. The grain size distribution of the Fe-8Al-4Ti alloy can exhibit a mode at approximately 11 pm, which can be substantially finer than the grain size of the Fe-6Al-6Ti composition. The reduction in grain size from the Fe-6Al-6Ti composition to the Fe-8Al-4Ti composition can be atributed to the B2 phase inducing a grain-pinning effect during recrystallization.

[0149] FIG. 9C corresponds to the Fe-10Al-2Ti alloy composition and shows an increased volume fraction of bcc phase 104 distributed throughout the fee matrix 102 in an elongated morphology. The Fe-10Al-2Ti composition can exhibit the finest grain size among the alloy compositions of the present disclosure, with a grain size distribution having a mode at approximately 7.5 pm. The fine grain size of the Fe-10Al-2Ti composition can be atributed to the B2 phase providing a grain-pinning effect during recrystallization, where the B2 grains can impede the migration of fee grain boundaries and can limit grain growth during the recrystallization treatment. The grain-pinning effect can result in finer grain sizes in dualphase alloys compared to single-phase fee alloys such as the Fe-6Al-6Ti composition.

[0150] FIG. 9D corresponds to the Fe-12A1 alloy composition and can exhibit a further increased bcc phase fraction with larger and more prominent bcc regions 104 dispersed within the fee matrix 102. The Fe-12A1 composition can exhibit a grain size distribution having a mode at approximately 11 pm. The grain size of the Fe-12A1 composition can be larger than the grain size of the Fe-10Al-2Ti composition despite the higher B2 phase fraction, which can be atributed to differences in the distribution and morphology of the B2 phase that can affect the grain-pinning efficiency during recrystallization.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0151] The alloy compositions of the present disclosure can exhibit unimodal grain size distributions for both fee grains and B2 grains. The unimodal grain size distribution can indicate that the recrystallization treatment produces a homogeneous grain structure without bimodal or multimodal grain size populations. The homogeneous grain structure can provide consistent mechanical properties throughout the alloy and can reduce the likelihood of strain localization at regions of grain size heterogeneity during deformation. The grain sizes of the fee grains and the B2 grains can vary by composition, with the Fe-10Al-2Ti composition exhibiting the finest grain size and the Fe-6Al-6Ti composition exhibiting the coarsest grain size among the alloy compositions characterized in FIGS. 9A-9D.

[0152] Referring to FIGS. 10A-10D, schematic diagrams illustrate different deformation modes operating in the alloy compositions of the present disclosure during uniaxial tensile loading at room temperature. As mentioned above, the fee grains 102 of the alloy compositions can exhibit deformation mechanisms including wavy slip, planar slip, high-density dislocation walls (HDDWs), and microbands (MBs) at different strain levels. The sequential activation of these deformation mechanisms can contribute to the multiple stages of strain hardening discussed with reference to FIGS. 7A-7D and FIG. 8. The alloy compositions of the present disclosure can exhibit medium to high stacking fault energy, which can prevent deformation twinning and deformation-induced martensitic transformation during tensile loading.

[0153] FIG. 10A shows wavy slip, also referred to as non-planar slip. During wavy slip, dislocations can glide on different crystallographic slip systems with a tendency to cross-slip, changing the slip systems and resulting in a wavy nature of slip. As illustrated in FIG. 10A, dislocations 106 during wavy slip can be distributed in a scattered, random pattern throughout the deforming region rather than being confined to specific crystallographic planes. The wavy slip mechanism can occur when dislocations are not confined to specific {111} planes and instead exhibit a tendency to cross-slip among multiple non-coplanar {111} planes. The wavy slip arrangement can be characteristic of alloys having medium to high stacking fault energy, where the extended dislocation width is sufficiently narrow to permit cross-slip between slip planes.

[0154] With continued reference to FIGS. 10A-10D, FIG. 10B shows planar slip. During planar slip, dislocations 108 can be confined to the primary slip system with the highest Schmid factor and can be unable to cross-slip to other slip systems. As illustrated in FIG. 10B, dislocations 108 during planar slip can be confined along parallel diagonal linesAttorney Docket No.: MIT 26210 PCT | 88212-432499representing the primary slip system. The planar slip mechanism can demonstrate a planar dislocation slip mechanism where destruction of the ordered structure via dislocation shearing facilitates easier dislocation glide on the primary slip system. The destruction of the ordered structure can be similar to the glide plane softening phenomenon observed in ordered intermetallic compounds. Planar slip can occur when dislocations shear through LI 2 precipitates, creating anti-phase boundaries that reduce the resistance to subsequent dislocation motion on the same slip plane.

[0155] FIG. 10C shows high-density dislocation walls (HDDWs) 110. The increased density of in-plane dislocations confined to the primary systems can produce long-range back stress, leading to the activation of other slip systems. Dislocations from other slip systems can cross-slip and join the primary slip system, forming directional dislocation arrays or bands. As illustrated in FIG. 10C, HDDWs 110 can appear as dense parallel formations of dislocations arranged in directional arrays aligned along crystallographic planes. The formation of HDDWs 110 can represent a transition from planar slip to a more complex dislocation structure where multiple slip systems become active due to the back stress generated by dislocation pileups on the primary slip system.

[0156] With continued reference to FIGS. 10A-10D, FIG. 10D shows microbands (MBs) 112. A further increase in dislocation density in the HDDWs, along with increased back-stress, can cause significant dislocation activity in other slip systems. At this stage, the activation of the cross-slip mechanism in the interwall spaces can cause dislocations to join secondary and tertiary slip systems, resulting in microband formation. As illustrated in FIG.10D, microbands 112 can appear as a dense cross-hatched pattern of intersecting dislocation bands where dislocations from multiple slip systems interact. The microbands can provide long-range back stress that can sustain the strain hardening rate during the later stages of deformation, thereby delaying plastic instability and enabling the alloy to achieve high uniform elongation.

[0157] The evolution of deformation mechanisms from wary slip to planar slip to HDDWs to microbands can occur progressively with increasing plastic strain during tensile loading. The transition between deformation mechanisms can depend on the local stress state, the precipitate distribution, and the accumulated dislocation density' within individual grains. Grains initially deforming by planar glide can evolve into HDDWs with increasing strain, whereas grains deforming by waxy slip can predominantly develop dislocation tangles. The coexistence of different deformation mechanisms within the same sample at aAttorney Docket No.: MIT 26210 PCT | 88212-432499given strain level can be attributed to variations in grain orientation and local microstructural features that affect the operative slip systems and the tendency for cross-slip.

[0158] The medium to high stacking fault energy of the alloy compositions of the present disclosure can prevent deformation twinning during tensile loading. Deformation twinning can occur in alloys having low stacking fault energy, where the extended dislocation width is sufficiently large to permit the nucleation and propagation of twin boundaries. The absence of deformation twinning in the alloy compositions of the present disclosure can be attributed to the stacking fault energy’ being above the threshold for twin nucleation. The medium to high stacking fault energy can also prevent deformation-induced martensitic transformation, which can occur in metastable austenitic alloys where the driving force for transformation exceeds the energy barrier for martensite nucleation. The absence of deformation-induced martensitic transformation can enable the alloy compositions to achieve high ductility without the premature damage nucleation that can be associated with phase transformation during deformation.

[0159] Referring to FIGS. 11A-14B, electron channeling contrast imaging (ECCI) micrographs and inverse pole figure (IPF) maps illustrate the evolution of deformation mechanisms at various plastic strain levels in the Fe-10Al-2Ti alloy composition. FIGS. 11A-14B demonstrate hoyv the deformation mechanisms transition from planar slip 106 at low strain levels to high-density dislocation walls (HDDWs) 110 at intermediate strain levels to microbands 112 at high strain levels, corresponding to the strain hardening stages discussed with reference to FIGS. 7A-7D and FIG. 8. The ECCI micrographs and IPF maps in FIGS.11A-14B further demonstrate that the deformation mechanisms show no strong crystallographic dependence on grain orientation.

[0160] At a plastic strain of approximately 5 percent, corresponding to Stage 2 of the strain hardening behavior, the Fe-10Al-2Ti alloy can exhibit planar slip as the dominant deformation mechanism within the fee grains. As shown in FIGS. 11 A-14B, ECCI micrographs of the Fe-10Al-2Ti alloy deformed to 5 percent plastic strain reveal dislocations confined to specific cry stallographic planes, characteristic of planar slip arising from glideplane softening. The planar slip features 106 can be indicated by the arrows in the ECCI micrographs in FIG. 11 A, showing dislocations aligned along {111} planes yvithin the fee grains 102. The confinement of dislocations to the primary7slip system during Stage 2 can be attributed to the shearing of LE precipitates by dislocations, which can create anti-phase boundaries that reduce the resistance to subsequent dislocation motion on the same slip plane.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0161] With continued reference to FIGS. 11 A-14B, the ECCI micrographs of the Fe-lOAl-2Ti alloy deformed to 5 percent plastic strain also reveal the presence of HDDWs 110 in some grains. The coexistence of planar slip 106 and HDDWs 110 at 5 percent plastic strain (0.05) can indicate that the transition from planar slip 106 to HDDWs 110 can begin during Stage 2 in some grains while other grains continue to deform by planar slip. The IPF map corresponding to the 5 percent plastic strain condition shows the distribution of deformation mechanisms as a function of grain orientation, where the red dots 106 indicate planar slip and green dots indicate non-planar slip 108. The distribution of these dots across the IPF map demonstrates that both planar slip 106 and non-planar slip 108 can occur across various grain orientations without a strong crystallographic preference for either deformation mechanism.

[0162] At a plastic strain of approximately 15 percent, corresponding to Stage 3 of the strain hardening behavior, the Fe-10Al-2Ti alloy can exhibit HDDWs as a prevalent deformation mechanism. As illustrated in FIGS. 11A-14B, ECCI micrographs of the Fe-10Al-2Ti alloy deformed to 15 percent plastic strain (0.15) reveal the formation of directional dislocation arrays aligned along crystallographic planes. The HDDWs 110 can be indicated by the arrows in the ECCI micrographs in FIG. 12B, for example, showing dense parallel formations of dislocations that have evolved from the planar slip structures observed at lower strain levels. The formation of HDDWs 110 during Stage 3 can be attributed to the increased density of in-plane dislocations that produce long-range back stress, leading to the activation of other slip systems and the cross-slip of dislocations to join the primary slip system.

[0163] The ECCI micrographs of the Fe-10Al-2Ti alloy deformed to 15 percent plastic strain (0.15) also reveal the formation of microbands 112 in some grains. The microbands 112 can be indicated by the arrows in the ECCI micrographs in FIGS. 12A-12B. showing dense cross-hatched patterns of intersecting dislocation bands. The coexistence of HDDWs 110 and microbands 112 at 15 percent plastic strain can indicate that the transition from HDDWs 110 to microbands 112 can begin during Stage 3 in some grains while other grains continue to exhibit HDDWs as the dominant deformation structure. The IPF map corresponding to the 15 percent plastic strain condition shows predominantly purple dots indicating HDDWs 110 distributed across various grain orientations, with some dots indicating non-planar slip 108. The distribution of deformation mechanisms across the IPF map demonstrates that HDDWs can form in grains having various crystallographic orientations without a strong preference for specific orientations.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0164] At a plastic strain of approximately 26 percent (0.26), corresponding to Stage 4 of the strain hardening behavior, the Fe-10Al-2Ti alloy can exhibit microbands as the dominant deformation mechanism. As shown in FIGS. 13A-13B, ECO micrographs of the Fe-lOAl-2Ti alloy deformed to 26 percent plastic strain reveal the formation of dense microband structures aligned along crystallographic planes. The microbands 112 can be indicated by the arrows in the ECCI micrographs of FIGS. 13A-13B, showing the characteristic cross-hatched pattern where dislocations from multiple slip systems interact to form dense dislocation bands. The formation of microbands 112 during Stage 4 can be attributed to the further increase in dislocation density in the HDDWs 110 and the increased back-stress that causes significant dislocation activity in secondary and tertiary slip systems.

[0165] With continued reference to FIGS. 11 A-14B, the IPF map corresponding to the 26 percent plastic strain (0.26) condition shows dots representing HDDWs 110 and dots representing microbands 112 distributed across the crystallographic orientation space. The distribution of purple and blue dots across the IPF map demonstrates that both HDDWs 110 and microbands 112 can form in grains having various crystallographic orientations. The absence of a strong clustering of specific deformation mechanisms at particular orientations in the IPF map indicates that the deformation mechanisms in the Fe-10Al-2Ti alloy show no strong crystallographic dependence on grain orientation. The grain orientation independence of the deformation mechanisms can enable homogeneous deformation throughout the poly crystalline microstructure, which can reduce the likelihood of strain localization and premature failure at grain boundaries.

[0166] At the fracture strain of approximately 39.5 percent (0.395). the Fe-10Al-2Ti alloy can continue to exhibit microbands 112 as the dominant deformation mechanism. As illustrated in FIGS. 14A-14B, ECCI micrographs of the Fe-10Al-2Ti alloy deformed to fracture reveal the continued presence of microband structures within the fee grains. The microbands 112 at fracture strain can be indicated by THE arrows in the ECCI micrographs of FIG. 14A, showing that the microband deformation mechanism can persist throughout Stage 4 until the onset of fracture. The IPF map corresponding to the fracture strain condition shows blue dots representing microbands distributed throughout thecry stallographic orientation space, confirming that microband formation can occur across various grain orientations at the fracture strain.

[0167] The evolution of deformation mechanisms from planar slip 106 at 5 percent strain to HDDWs 110 at 15 percent strain to microbands 112 at 26 percent strain and beyond canAttorney Docket No.: MIT 26210 PCT | 88212-432499correspond to the transitions between strain hardening stages observed in the strain hardening rate versus true strain plots discussed with reference to FIGS. 7A-7D and FIG. 8. The transition from planar slip 106 to HDDWs 110 can correspond to the transition from Stage 2 to Stage 3, where the strain hardening contribution from LI 2 precipitate shearing begins to saturate and the formation of directional dislocation arrays provides continued strain hardening. The transition from HDDWs 110 to microbands 112 can correspond to the transition from Stage 3 to Stage 4. where the B2s within the fee grains promote cross-slip and microband formation that sustains the strain hardening rate during the later stages of deformation.

[0168] The absence of a strong crystallographic dependence of deformation mechanisms on grain orientation in the Fe-10Al-2Ti alloy can be attributed to the hierarchical multiprecipitate microstructure that provides multiple mechanisms for dislocation storage and strain hardening. The Lb precipitates, B2s, and bcc precipitates distributed throughout the fee and B2 grains can interact with dislocations regardless of the grain orientation, enabling similar deformation mechanisms to operate across grains having different crystallographic orientations. The grain orientation independence of the deformation mechanisms can contribute to the high uniform elongation achieved by the Fe-10Al-2Ti alloy by enabling homogeneous strain accommodation throughout the poly crystalline microstructure without preferential strain localization in grains having specific orientations.

[0169] Referring to FIG. 15, a scanning electron microscope (SEM) micrograph illustrates the hierarchical microstructure of the aged Fe-10Al-2Ti alloy at a scale of 1 micrometer. The SEM micrograph displays the dual-phase matrix architecture of the alloy composition, showing the distribution of different grain tvpes and precipitate phases that contribute to the mechanical properties of the material. Three distinct regions are identified and marked with dashed rectangular boxes labeled A, B, and C in FIG. 15, each corresponding to a different microstructural constituent within the hierarchical architecture.

[0170] Region A in FIG. 15 can correspond to an area within a face-centered cubic (fee) grain 102 containing LL precipitates 114 (as shown in FIGS. 16B-16C). The fee grains 102 can serve as one of the two matrix phases in the dual-phase architecture of the Fe-10Al-2Ti alloy. The LE precipitates 114 within the fee grains can be distributed at the nanoscale and can provide strengthening through dislocation-precipitate interactions during deformation. The LI2 precipitates 114 can govern Stage 2 of the strain hardening behavior, as discussed with reference to FIG. 8, where dislocations can interact with the LEAttorney Docket No.: MIT 26210 PCT | 88212-432499precipitates 114 through shearing mechanisms that create anti-phase boundaries within the precipitates.

[0171] With continued reference to FIG. 15, Region B can identify aB2 grain containing body-centered cubic (bcc) precipitates 104 (as shown in FIG. 17B). The B2 grains can serve as the second matrix phase in the dual-phase architecture of the Fe-10Al-2Ti alloy. The bcc precipitates 104 within the B2 grains can be visible as cuboidal features within the B2 matrix in FIG. 15. The B2 grains containing bcc precipitates 104 can contribute to Stage 3 of the strain hardening behavior, where the B2 grains can accommodate plastic strain through dislocation activity within the B2 matrix and through shearing of the bcc precipitates. The Fe-10Al-2Ti alloy can contain approximately 6.5% B2 matrix phase distributed throughout the microstructure.

[0172] Region C in FIG. 15 indicates a smaller feature representing B2 and bcc phases present w ithin the fee grains. The presence of B2 + bcc phases within the fee grains 102 can represent a third microstructural constituent that distinguishes the Fe-10Al-2Ti alloy from single-precipitate alloy systems. The B2 precipitates within the fee grains can control Stage 4 of the strain hardening behavior through microband formation, as discussed with reference to FIG. 8. The Fe-10Al-2Ti alloy can contain approximately 3.5% bcc precipitates distributed within the microstructure.

[0173] The hierarchical microstructure shown in FIG. 15 can enable the superior mechanical properties of the Fe-10Al-2Ti alloy by providing multiple mechanisms for strain hardening at different strain levels. The dual-phase matrix architecture composed of fee grains and B2 grains can provide two distinct matrix phases that can accommodate plastic strain through different deformation mechanisms. The multiple precipitate ty pes distributed within the dual-phase matrix, including Lh precipitates in the fee grains (dashed box A), bcc precipitates in the B2 grains (dashed box B), and B2 + bcc phases within the fee grains (dashed box C), can provide sequential activation of strain hardening mechanisms that sustain work hardening throughout the deformation process. The combination of the dual-phase matrix with the multiple precipitate types can enable the Fe-10Al-2Ti alloy to achieve the four distinct strain hardening stages discussed with reference to FIGS. 7A-7D and FIG. 8, thereby enabling the alloy to achieve both high strength and high ductility without relying on martensite or metastable austenite phases.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0174] Referring to FIGS. 16A-16D, transmission electron microscopy (TEM) characterization illustrates the microstructure of fee grains containing LE precipitates in the aged Fe-10Al-2Ti alloy subjected to heat treatment at 650 degrees Celsius for 15 hours. FIG.1 A shows the aged microstructure of the Fe-10Al-2Ti alloy at a scale of 1 micrometer, revealing the grain structure with a region of interest indicated by the dashed box A. The microstructure displayed in FIG. 16A exhibits multiple grains with vary ing contrast, indicating different crystallographic orientations and phase distributions within the aged alloy. Dark regions visible within the microstructure of FIG. 16A correspond to voids or pores, while bright linear features along grain boundaries indicate the presence of secondary7phases or precipitates.

[0175] With continued reference to FIGS. 16A-16D, FIG. 16B presents an electron diffraction pattern taken along the

[0011] zone axis from an fee grain of the aged Fe-10Al-2Ti alloy. The electron diffraction pattern in FIG. 16B shows LE superlattice spots in addition to the fundamental fee reflections. The presence of LE superlattice spots in the electron diffraction pattern confirms the presence of L L-ordered precipitates within the fee matrix phase. The LE superlattice reflections arise from the ordered arrangement of atoms in the Lb structure, where the ordering creates additional diffraction spots at positions that are forbidden in the disordered fee structure.

[0176] FIG. 16C displays an electron diffraction pattern along the

[0112] zone axis from an fee grain of the aged Fe-10Al-2Ti alloy. The electron diffraction pattern in FIG. 16C also exhibits LE superlattice spots that further verify the ordered precipitate structure within the fee matrix. The observation of LE superlattice reflections along multiple zone axes, including the

[0011] zone axis shown in FIG. 16B and the

[0112] zone axis shown in FIG. 16C, confirms that the Lb ordering is present throughout the precipitates rather than being limited to specific crystallographic orientations. The electron diffraction patterns in FIGS. 16B and 16C demonstrate that the fee grains 102 of the aged Fe-10Al-2Ti alloy contain LE-ordered precipitates that contribute to the strengthening of the alloy through dislocation-precipitate interactions during deformation.

[0177] FIG. 16D shows a dark-field TEM image obtained using a diffraction vector g equal to 100, revealing a uniform distribution of fine precipitates throughout the fee matrix 102 at a nanometer scale. The precipitates in FIG. 1 D appear as bright contrast features dispersed homogeneously within the darker matrix background. The uniform distribution of LE precipitates shown in FIG. 16D demonstrates that the aging heat treatment produces aAttorney Docket No.: MIT 26210 PCT | 88212-432499homogeneous precipitate microstructure within the fee grains 102. The Lh precipitates within the fee matrix 102 can have an average size and are uniformly distributed throughout the fee grains 102, which can provide consistent strengthening throughout the fee phase during deformation.

[0178] The LI 2 precipitates within the fee matrix can exhibit a lattice misfit of approximately -0.84% between the LL and fee phases. The negative lattice misfit indicates that the lattice parameter of the LL precipitates can be smaller than the lattice parameter of the fee matrix. The lattice misfit of approximately -0.84% can indicate coherency between the LL precipitates and the fee matrix, where the precipitate-matrix interface maintains cry stallographic continuity7without the formation of misfit dislocations. The coherent nature of the LI2 precipitates can enable effective strengthening through coherency strain hardening and through the creation of anti-phase boundaries when dislocations shear through the precipitates during deformation.

[0179] The coherent LL precipitates within the fee matrix 102 can contribute to the strain hardening behavior of the alloy compositions of the present disclosure. The coherency between the LL precipitates and the fee matrix can enable dislocations to shear through the precipitates during Stage 2 of the strain hardening behavior, creating anti-phase boundaries that provide resistance to subsequent dislocation motion. The uniform distribution of LL precipitates shown in FIG. 16D can ensure that dislocations encounter precipitates throughout the fee grains 102 during deformation, providing consistent strain hardening throughout the fee phase. The TEM characterization results shown in FIGS. 16A-16D confirm the hierarchical microstructure achieved through the aging heat treatment, which contributes to the mechanical properties of the alloy system.

[0180] The alloy compositions of the present disclosure can include aluminum (Al) approximately in a range of about 5 at.% to about 14 at.%. In some embodiments, the Al can be approximately in a range of about 7 at.% to about 10 at.%. The Fe-10Al-2Ti composition characterized in FIGS. 16A-16D can contain Al at approximately 10 at.%, which falls within the range of about 7 at.% to about 10 at.%. The Al content within the range of about 7 at.% to about 10 at.% can provide a balance between LL precipitate stability7and B2 phase formation, enabling the hierarchical multi-precipitate microstructure that provides the multiple strain hardening stages discussed with reference to FIGS. 7A-7D and FIG. 8.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0181] A method of forming an alloy of the present disclosure can include providing Al approximately in a range of about 5 at.% to about 14 at.%. In some embodiments, the method can include providing Al approximately in a range of about 7 at.% to about 10 at.%. The Al content within the range of about 7 at.% to about 10 at.% can promote the formation of Lh precipitates within the fee grains during the aging treatment, as demonstrated by the TEM characterization results shown in FIGS. 16A-16D. The Al content can also influence the volume fraction of the B2 phase, as discussed with reference to FIGS.1 A-1E, where increasing Al content can promote the formation of the B2 phase at the expense of the fee phase. The selection of Al content within the range of about 7 at.% to about 10 at.% can enable the formation of the dual-phase multi -precipitate microstructure that provides the superior strength-ductility combination of the alloy compositions of the present disclosure.

[0182] Referring to FIGS. 17A-17D, transmission electron microscopy (TEM) characterization illustrates the microstructure of B2 grains containing body-centered cubic (bcc) precipitates 104 in the aged Fe-10Al-2Ti alloy subjected to heat treatment at 650 degrees Celsius for 15 hours. FIG. 17A shows a TEM micrograph of the aged microstructure with a scale bar of 1 micrometer, revealing the grain structure with dark regions corresponding to voids or pores and a central region highlighted by a dashed box B indicating an area of interest for higher magnification analysis. The microstructure exhibited in FIG. 17A includes characteristic features such as linear striations and contrast variations indicative of the hierarchical phase distribution within the alloy. The B2 grains shown in FIG. 17A can serve as the second matrix phase in the dual-phase architecture of the Fe-lOAl-2Ti alloy, complementing the fee grains 102 containing LE precipitates discussed with reference to FIGS. 16A-16D.

[0183] With continued reference to FIGS. 17A-17D, FIG. 17B presents a selected area electron diffraction pattern along the

[0001] zone axis from a B2 grain of the aged Fe-lOAl-2Ti alloy. The electron diffraction pattern in FIG. 17B displays B2 superlattice spots marked by arrows, which confirm the presence of B2 ordering in the material. The diffraction pattern shows the fundamental reflections along with the superlattice reflections characteristic of the ordered B2 phase. The B2 superlattice reflections arise from the ordered arrangement of atoms in the B2 structure, where the CsCI-lype ordering creates additional diffraction spots at positions that are forbidden in the disordered bcc structure. The presence of B2 superlatticeAttorney Docket No.: MIT 26210 PCT | 88212-432499spots in the electron diffraction pattern confirms that the B2 grains of the aged Fe-10Al-2Ti alloy contain an ordered B2 matrix phase with disordered bcc precipitates distributed therein.

[0184] FIG. 17C provides a higher magnification view corresponding to the boxed region B in FIG. 17A. showing the detailed microstructural features including cuboidal precipitates distributed within the B2 matrix phase. The image in FIG. 17C reveals the morphology and distribution of the bcc precipitates within the B2 matrix, with the 010 direction indicated by an arrow. The bcc precipitates within the B2 matrix can be cuboidal in shape, as illustrated in FIG. 17C, where the precipitates exhibit well-defined facets aligned along crystallographic directions of the B2 matrix. The cuboidal morphology of the bcc precipitates 104 can be attributed to the elastic anisotropy of the B2 matrix and the coherent nature of the precipitatematrix interface, which can favor the development of faceted precipitate shapes that minimize the total interfacial and elastic strain energy.

[0185] With continued reference to FIGS. 17A-17D, FIG. 17D presents a dark-field TEM image obtained using the g-vector equal to 010 reflection, with a scale bar of 200 nanometers, showing the distribution of precipitates and dislocation structures within the aged alloy. The image in FIG. 17D reveals the uniform distribution of cuboidal bcc precipitates within the B2 ordered matrix, along with linear features corresponding to dislocation arrays aligned along crystallographic directions. The contrast in FIG. 17D highlights the ordered B2 matrix phase while the darker cuboidal regions correspond to the disordered bcc precipitates. The darkfield imaging conditions used in FIG. 17D enable the distinction between the ordered B2 matrix and the disordered bcc precipitates based on the diffraction contrast arising from the different crystal structures.

[0186] The bcc precipitates 104 within the B2 matrix can have an average size of approximately 18.5 ± 5.5 nm. The average precipitate size of approximately 18.5 nm can be determined from measurements of at least 500 precipitates from different grains using bright-field and dark-field TEM images. The precipitate size calculation can be based on areal equivalent circle radius, where the edge lengths of the cuboidal precipitates can be measured and the area of each precipitate can be equated with the equivalent area of a circle of radius r. The standard error of ± 5.5 nm can represent the standard deviation of the average precipitate size across the measured precipitate population. The uniform distribution and consistent size of the bcc precipitates within the B2 matrix, as shown in FIG. 17D, can provide consistent strengthening throughout the B2 phase during deformation.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0187] The bcc precipitates within the B2 matrix can exhibit a lattice misfit of approximately -0.26% between the bcc and B2 phases. The lattice misfit can be calculated using the formula:

[0188] where aPis the lattice parameter of the precipitate phase and aM is the lattice parameter of the matnx phase. The negative lattice misfit of approximately -0.26% indicates that the lattice parameter of the bcc precipitates is smaller than the lattice parameter of the B2 matrix. The lattice misfit of approximately -0.26% can be smaller in magnitude than the lattice misfit of approximately -0.84% between the Lb precipitates and the fee matrix discussed with reference to FIGS. 16A-16D. The smaller lattice misfit between the bcc precipitates and the B2 matrix can indicate a higher degree of coherency at the precipitatematrix interface compared to the Lb / fcc interface.

[0189] The lattice misfit of approximately -0.26% between the bcc and B2 phases can indicate coherency between the bcc precipitates and the B2 matrix, where the precipitatematrix interface maintains crystallographic continuity without the formation of misfit dislocations. The coherent nature of the bcc precipitates within the B2 matrix can enable effective strengthening through coherency strain hardening and through dislocationprecipitate interactions during deformation. The B2 grains containing coherent bcc precipitates can contribute to Stage 3 of the strain hardening behavior, where the B2 grains can accommodate plastic strain through dislocation activity within the B2 matrix and through shearing of the bcc precipitates.

[0190] The alloy compositions of the present disclosure can include tantalum (Ta) at about 2 at.%. The Ta content of about 2 at.% can be selected to maximize the y'-stabilizing effect of Ta while avoiding the formation of deleterious topologically close-packed (TCP) phases. In some embodiments, the Ta concentration can be fixed at about 2 at.% because the solubility of Ta in Fe is limited to less than 1 at.%, although the presence of Ni and Co can increase Ta solubility. Additions of Ta beyond 2 at.% in CoNi-based alloys can promote TCP phase formation, and therefore the Ta content can be limited to about 2 at.% in the alloy compositions of the present disclosure. The Fe-10Al-2Ti composition characterized in FIGS.17A-17D contains Ta at approximately 2 at.%, which can contribute to the stability of the LG precipitates within the fee grains and can influence the partitioning of elements between the fee and B2 phases.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0191] A method of forming an alloy of the present disclosure can include providing Ta at about 2 at.%. The method can include combining the Ta with the other alloying elements including Fe, Ni, Co, Al, Ti, V, and B to form the alloy system. The Ta content of about 2 at.% can be provided using pure Ta metal during the arc-melting process, where the Ta can be incorporated into the alloy melt along with the other constituent elements. The Ta can partition between the fee and B2 phases during solidification and subsequent heat treatment, with a distribution coefficient of approximately 0.95 indicating that Ta can partition slightly to the B2 phase relative to the fee phase.

[0192] A method of forming an alloy of the present disclosure can include providing Ti at about 2 at.%. The method can include combining the Ti with the other alloying elements including Fe, Ni, Co, Al, Ta, V, and B to form the alloy system. The Ti content of about 2 at.% can be provided using pure Ti metal during the arc-melting process, where the Ti can be incorporated into the alloy melt along with the other constituent elements. That is, the Ti content of about 2 at.% can be selected to utilize the y'-stabilizing effect of Ti without compromising mechanical performance. An increase in Ti content can reduce the shearability' of the Lh precipitates due to the higher planar fault energy of the precipitates, which can inhibit planar slip and can lead to the formation of dislocation walls and microbands that can degrade the mechanical properties. The Ti can partition between the fee and B2 phases during solidification and subsequent heat treatment, with a distribution coefficient of approximately 0.61 indicating that Ti can partition to the B2 phase relative to the fee phase. The partitioning of Ti to the B2 phase can influence the composition and properties of the bcc precipitates within the B2 matrix, as characterized in FIGS. 17A-17D.

[0193] Referring to FIG. 18. an X-ray diffraction pattern illustrates the crystallographic phases present in the Fe-10Al-2Ti alloy composition of the present disclosure. The horizontal axis of FIG. 18 represents the diffraction angle 29 measured in degrees, ranging from approximately 1 to 7 degrees. The vertical axis of FIG. 18 represents intensity measured in arbitrary’ units. The diffraction pattern displayed in FIG. 18 exhibits multiple peaks at various 29 positions, with each peak labeled to indicate the corresponding crystallographic phase and Miller indices. The peaks in FIG. 18 are identified as belonging to several distinct phases including FCC-Lb and BCC-B2 phases and their associated crystallographic planes.

[0194] The X-ray diffraction pattern of the Fe-10Al-2Ti alloy shown in FIG. 18 confirms the presence of both fcc-based and bcc-based phases within the alloy microstructure. TheAttorney Docket No.: MIT 26210 PCT | 88212-432499strongest peak in FIG. 18 appears at approximately 3.3 degrees and corresponds to the FCC-L (111) reflection. Additional FCC-Lh peaks are observed at the (200), (220), (311), and (222) reflections. A peak labeled as LL (100) appears at approximately 2 degrees, which corresponds to a superlattice reflection arising from the ordered LL structure. The presence of the LG (100) superlattice reflection in the X-ray diffraction pattern confirms the presence of LL-ordered precipitates within the fee matrix, consistent with the TEM characterization results discussed with reference to FIGS. 16A-16D.

[0195] With continued reference to FIG. 18, peaks corresponding to the BCC-B2 phase are identified at the (100), (110), (200), (211), and (220) reflections. The (100) reflection of the B2 phase corresponds to a superlattice reflection arising from the ordered B2 structure, which confirms the presence of B2 ordering within the bcc-based grains. The presence of both FCC-LI2 and BCC-B2 phase peaks in the diffraction pattern of FIG. 18 confirms the dualphase microstructure of the aged Fe-I0Al-2Ti alloy, demonstrating the coexistence of facecentered cubic and body -centered cubic ordered phases within the material. The X-ray diffraction characterization results shown in FIG. 18 corroborate the TEM observations discussed with reference to FIGS. 16A-16D and FIGS. 17A-17D, which identified LI2 precipitates 114 within fee grains 102 and bcc precipitates 104 within B2 grains.

[0196] The X-ray diffraction pattern of the Fe-10Al-2Ti alloy shown in FIG. 18 does not exhibit peaks corresponding to deformation-induced martensite or other metastable phases. The absence of martensite peaks in the diffraction pattern confirms that the Fe-lOAl-2Ti alloy composition achieves the dual-phase multi-precipitate microstructure w ithout relying on martensitic transformation. The alloy compositions of the present disclosure can thus achieve high strength and high ductility through the hierarchical multi -precipitate microstructure rather than through deformation-induced phase transformation mechanisms that can be associated with premature damage nucleation.

[0197] The alloy compositions of the present disclosure can include vanadium (V) at about 5 at.%. The V content of about 5 at.% can be selected to stabilize both the bcc phase and the LL precipitates within the alloy microstructure. V can be an a-BCC stabilizer in the Fe system and can stabilize LI2 precipitates through selective partitioning to the y' precipitates, similar to the behavior of V in Ni-based and Co-based superalloys. The Fe-10Al-2Ti composition characterized in FIG. 18 contains V at approximately 5 at.%, which can contribute to the stability of both the bcc phase within the B2 grains and the LL precipitates within the fee grains. The distribution coefficient of V between the fee and B2 phases can beAttorney Docket No.: MIT 26210 PCT | 88212-432499greater than 1, indicating that V can partition preferentially to the fee phase relative to the B2 phase. The preferential partitioning of V to the fee phase can enhance the stability of the LI 2 precipitates within the fee grains while also contributing to bee phase stabilization through the overall alloy composition.

[0198] A method of forming an alloy of the present disclosure can include providing V at about 5 at.%. The method can include combining the V with the other alloying elements including Fe, Ni, Co, Al, Ta, Ti, and B to form the alloy system. The V content of about 5 at.% can be provided using pure V metal during the arc-melting process, where the V can be incorporated into the alloy melt along with the other constituent elements. The V can partition between the fee and B2 phases during solidification and subsequent heat treatment, contributing to the phase stability7and precipitate distribution that enables the hierarchical multi-precipitate microstructure characterized by X-ray diffraction in FIGS. 18 and 20.

[0199] The alloy compositions of the present disclosure can include boron (B) at about 0.02 at.% to 0.3 at%. B can be added to enhance grain boundary cohesion, although caution is encouraged to avoid promoting the formation of deleterious boride phases at grain boundaries. The B content of about 0.02 at.% can be selected to enhance grain boundary7cohesion without promoting the formation of deleterious boride phases at grain boundaries. B additions can enhance mechanical properties in Fe-based, Ni-based. and Cobased alloys through chemical bonding between B and neighboring metal atoms. The B-metal hybridization at grain boundaries can fill the metal d-band, thereby' restoring cohesive strength comparable to bulk metal-metal bonds. The relatively low electronegativity7of B can prevent depletion of metal-metal bonding electrons, and B can form covalent bonds with surrounding metal atoms that enhance cohesion. The B content can be limited to about 0.02 at.% to avoid the formation of Ta-rich borides at grain boundaries that could degrade the mechanical properties of the alloy.

[0200] A method of forming an alloy of the present disclosure can include providing B at about 0.02 at.%. The method can include combining the B with the other alloying elements including Fe, Ni, Co, Al, Ta, Ti, and V to form the alloy system. The B content of about 0.02 at.% can be provided using a master alloy of iron and boron (Fe-B) during the arc-melting process to prevent boron evaporation during casting. The use of the Fe-B master alloy can ensure that the boron content is retained in the alloy during the high-temperature arc-melting process. The B can segregate to grain boundaries during solidification and subsequent heat treatment, where the B can enhance grain boundary cohesion. The enhanced grain boundaryAttorney Docket No.: MIT 26210 PCT | 88212-432499cohesion provided by the B content of about 0.02 at.% can contribute to the high ductility achieved by the alloy compositions of the present disclosure by reducing the likelihood of intergranular fracture during tensile deformation.[02011 FIG. 20 illustrates X-ray diffraction patterns that show the evolution of phase constitution with composition across the Fe-6Al-6Ti, Fe-8Al-4Ti, and Fe-12A1 alloy compositions of the present disclosure. The horizontal axis of FIG. 20 represents the diffraction angle 20 measured in degrees, ranging from approximately 3 degrees to 11 degrees. The vertical axis of FIG. 20 represents intensity measured in arbitrary units. Three diffraction patterns are displayed in FIG. 20, stacked vertically, with Fe-12A1 shown at the top, Fe-8Al-4Ti shown in the middle, and Fe-6Al-6Ti shown at the bottom. Multiple diffraction peaks are labeled throughout the patterns in FIG. 20, indicating the presence of various phases including FCC + Lb phases and BCC + B2 phases.

[0202] The X-ray diffraction pattern of the Fe-6Al-6Ti alloy shown at the bottom of FIG.20 exhibits peaks corresponding to the disordered fee phase and the ordered Lb phase. The Fe-6Al-6Ti diffraction pattern shows the fewest and smallest peaks among the three compositions displayed in FIG. 20, with peaks identified at positions corresponding to FCC + Lb crystallographic planes such as (100), (110), (111), (200), (220), (311), and (222). The absence of BCC + B2 phase peaks in the Fe-6Al-6Ti diffraction pattern confirms that the Fe-6Al-6Ti alloy exhibits a single-phase fcc-based microstructure with Lb precipitates, consistent with the EBSD phase map shown in FIG. 9A and the thermodynamic predictions discussed with reference to FIG. IB. The single-phase fee + LL microstructure of the Fe-6Al-6Ti alloy can be attributed to the VEC of approximately 8.0 that stabilizes the fee phase without promoting B2 phase formation.

[0203] With continued reference to FIG. 20, the X-ray diffraction pattern of the Fe-8Al-4Ti alloy shown in the middle of FIG. 20 exhibits additional peaks beyond those of fee and LL, which are indexed to the disordered bcc and ordered B2 phases. The Fe-8Al-4Ti diffraction pattern shows peaks corresponding to BCC + B2 phases at positions such as (110), (200), (211), and (220), in addition to the FCC + Lb peaks observed in the Fe-6Al-6Ti pattern. A peak corresponding to the B2 (100) superlattice reflection is also present in the Fe-8Al-4Ti diffraction pattern, confirming the presence of B2 ordering within the bcc-based grains. The presence of both FCC + Lb and BCC + B2 phase peaks in the Fe-8Al-4Ti diffraction pattern confirms the dual-phase microstructure of the aged Fe-8Al-4Ti alloy, consistent with theAttorney Docket No.: MIT 26210 PCT | 88212-432499EBSD phase map show n in FIG. 9B and the thermodynamic predictions discussed with reference to FIG. 1C.

[0204] The X-ray diffraction pattern of the Fe-12A1 alloy shown at the top of FIG. 20 exhibits the most prominent peaks with the greatest number of labeled phase identifications among the three compositions. The Fe-12A1 diffraction pattern shows peaks corresponding to FCC + Lb phases and BCC + B2 phases, similar to the Fe-8Al-4Ti pattern, but with increased intensity of the BCC + B2 phase peaks relative to the FCC + Lb phase peaks. The increased intensity of the BCC + B2 phase peaks in the Fe-12A1 diffraction pattern can be attributed to the higher volume fraction of the B2 phase in the Fe-12A1 alloy compared to the Fe-8Al-4Ti alloy, consistent with the EBSD phase map shown in FIG. 9D and the thermodynamic predictions discussed with reference to FIG. IE. The diffraction peaks in FIG. 20 are indexed to the fee. LI2, bcc, and B2 phases, confirming the presence of multiple crystallographic phases in the Fe-8Al-4Ti and Fe-12A1 alloy compositions.

[0205] The evolution of phase constitution with composition illustrated in FIG. 20 demonstrates that increasing aluminum content and decreasing titanium content can promote the formation of the B2 phase at the expense of the fee phase. The Fe-6Al-6Ti alloy having the lowest aluminum content exhibits only fee + LE phases, whereas the Fe-8Al-4Ti and Fe-12A1 alloys having higher aluminum contents exhibit fee + LE + bcc + B2 phases. The progression from single-phase fee + LE microstructure in the Fe-6Al-6Ti alloy to dual-phase fee + bcc + LE + B2 microstructure in the Fe-8Al-4Ti and Fe-12A1 alloys can be attributed to the decrease in VEC with increasing aluminum content, which can stabilize the fee + bcc phase field as discussed with reference to FIGS. 1A-1E.

[0206] The X-ray diffraction characterization results shown in FIGS. 18 and 20 confirm that the alloy compositions of the present disclosure can achieve the dual-phase multiprecipitate microstructure through compositional design based on VEC considerations. The XRD and TEM analyses demonstrate that the aged Fe-6Al-6Ti alloy contains only the fee matrix and LG-ordered precipitate phases, whereas the aged Fe-8Al-4Ti, Fe-10Al-2Ti, and Fe-12A1 alloys exhibit multiple phases including fee, LL, bcc, and B2. The presence of multiple crystallographic phases in the Fe-8Al-4Ti, Fe-10Al-2Ti. and Fe-12A1 alloys can enable the multiple stages of strain hardening discussed with reference to FIGS. 7A-7D and FIG. 8, where each phase and precipitate type can contribute to a distinct strain hardening mechanism during tensile deformation.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0207] Referring to FIGS. 19A-19D, atom probe tomography (APT) compositional analysis illustrates the elemental distribution in B2 grain regions and fee grain regions of the aged Fe-10Al-2Ti alloy. The APT characterization can provide nanoscale compositional information that reveals the partitioning behavior of alloying elements between the matrix phases and the precipitate phases within the hierarchical microstructure. FIGS. 19A and 19B correspond to a B2 grain containing bcc precipitates, while FIGS. 19C and 19D correspond to an fee grain containing Lh precipitates.

[0208] FIG. 19A shows a one-dimensional compositional profile along the Z-axis for a B2 grain with bcc precipitates. The vertical axis of FIG. 19A represents concentration in atomic percent, and the horizontal axis represents position in nanometers. Multiple elemental traces are displayed in FIG. 19A, showing periodic concentration variations that correspond to precipitate locations within the B2 matrix. The periodic variations in elemental concentration along the Z-axis demonstrate the alternating composition between the B2 matrix regions and the bcc precipitate regions. The bcc precipitates within the B2 matrix can be enriched in Fe, V, and Co, while the B2 matrix can be enriched in Ni. The enrichment of Fe, V, and Co in the bcc precipitates and the enrichment of Ni in the B2 matrix can be observed as anti-correlated concentration variations in the compositional profile of FIG. 19A, where peaks in Fe, V. and Co concentration correspond to troughs in Ni concentration at the precipitate locations.

[0209] With continued reference to FIGS. 19A-19D, FIG. 19B shows a three-dimensional reconstruction of the B2 grain region with bcc precipitates 114 rendered as ellipsoidal shapes distributed within the B2 matrix. The vertical axis of FIG. 19B represents the X-axis position in nanometers, and the horizontal axis represents the Z-axis position in nanometers. The arrow in FIG. 19B indicates specific precipitate features within a planar region 116 that represents a cross-section through the reconstructed volume. The three-dimensional reconstruction in FIG. 19B illustrates the spatial distribution of the bcc precipitates within the B2 matrix, showing that the precipitates can be uniformly distributed throughout the B2 grain. The ellipsoidal morphology of the bcc precipitates shown in FIG. 19B can be consistent with the cuboidal precipitate morphology observed in the TEM characterization discussed with reference to FIGS. 17A-17D, where the apparent ellipsoidal shape in the APT reconstruction can result from the projection of cuboidal precipitates along the analysis direction.

[0210] FIG. 19C shows a three-dimensional reconstruction of an fee grain region 102 with Lb precipitates 104. The vertical axis of FIG. 19C represents the X-axis position inAttorney Docket No.: MIT 26210 PCT | 88212-432499nanometers, and the horizontal axis represents the Z-axis position in nanometers. The arrow 104 in FIG. 19C highlights a specific precipitate cluster within the reconstructed volume. The three-dimensional reconstruction in FIG. 19C illustrates the spatial distribution of the LG precipitates 104 within the fee matrix 102, showing that the precipitates can be uniformly distributed throughout the fee grain. The uniform distribution of LG precipitates 104 shown in FIG. 19C can be consistent with the dark-field TEM observations discussed with reference to FIG. 16D, which demonstrated a homogeneous precipitate microstructure within the fee grains 102.

[0211] With continued reference to FIGS. 19A-19D, FIG. 19D shows a one-dimensional compositional profile along the Z-axis for the fee grain 102 with LL precipitates 104. The vertical axis of FIG. 19D represents concentration in atomic percent, and the horizontal axis represents position in nanometers. Multiple elemental traces are displayed in FIG.19D showing concentration variations that correspond to the periodic distribution of LI2 precipitates throughout the fee matrix. The LE precipitates can be enriched in Ni, Al, Ti, and Ta, while the fee matrix can retain Fe, Co, and V. The enrichment of Ni, Al, Ti, and Ta in the LG precipitates and the retention of Fe, Co, and V in the fee matrix can be observed as anti-correlated concentration variations in the compositional profile of FIG. 19D, where peaks in Ni, AL Ti, and Ta concentration correspond to troughs in Fe, Co, and V concentration at the precipitate locations.

[0212] The elemental partitioning behavior between the fee and B2 phases of the alloy compositions of the present disclosure can be quantified using distribution coefficients. Measuring composition with wavelength-dispersive spectroscopy (WDS) under the EPMA allows calculation of the distribution coefficient Ki for each element i, which can be defined as the ratio of the concentration of element i in the fee phase to the concentration of element i in the B2 phase. When measured in Fe-10Al-2Ti, Fe, Co, and V can preferentially partition to the fee phase, with distribution coefficients Ki greater than 1. The distribution coefficient of Fe can be approximately 1.47, the distribution coefficient of Co can be approximately 1.10, and the distribution coefficient of V can be approximately 1.38 in the Fe-10Al-2Ti alloy. The distribution coefficients greater than 1 for Fe, Co, and V indicate that these elements can be present at higher concentrations in the fee phase compared to the B2 phase, although it can be concluded that no strong preferential partitioning occurs between the two phases in all the elements.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0213] Ni, Al, Ta, and Ti can preferentially partition to the B2 phase, with distribution coefficients Ki less than 1. The distribution coefficient of Ni can be approximately 0.90. the distribution coefficient of Al can be approximately 0.44, the distribution coefficient of Ta can be approximately 0.95, and the distribution coefficient of Ti can be approximately 0.61 in the Fe-10Al-2Ti alloy. The distribution coefficients less than 1 forNi, Al, Ta, and Ti indicate that these elements can be present at higher concentrations in the B2 phase compared to the fee phase. The preferential partitioning of Ni and Al to the B2 phase can be consistent with the B2-NiAl compound stoichiometry, where Ni and Al can be the primary constituents of the ordered B2 structure.

[0214] The elemental partitioning behavior observed in the APT compositional analysis of FIGS. 19A-19D can influence the composition and properties of the precipitate phases within the hierarchical microstructure. The enrichment of Ni, Al, Ti, and Ta in the LL precipitates can be consistent with the y'-NisAl-type structure of the Lk phase, where Ni occupies the face-centered positions and Al, Ti, and Ta can substitute on the comer positions of the ordered structure. The retention of Fe, Co, and V in the fee matrix can maintain the disordered fee structure of the matrix phase while providing solid solution strengthening. The enrichment of Fe, V, and Co in the bcc precipitates within the B2 matrix can result in a disordered bcc structure for the precipitates, while the enrichment of Ni in the B2 matrix can maintain the ordered B2-NiAl-type structure of the matrix phase. The elemental partitioning behavior can thus contribute to the formation of the hierarchical multiprecipitate microstructure that provides the multiple strain hardening stages discussed with reference to FIGS. 7A-7D and FIG. 8.

[0215] Referring to FIG. 21A. a weak-beam dark-field TEM image illustrates the deformation substructure in fee grains of the Fe-10Al-2Ti alloy during Stage 2 of the strain hardening behavior. FIG. 21 A shows the microstructure of the alloy at a scale of 100 nm, with an inset displaying a selected area electron diffraction pattern along the

[0001] zone axis. The diffraction pattern in the inset (i) of FIG. 21 A shows spots characteristic of the Fe-10Al-2Ti composition, including both fundamental fee reflections and LE superlattice reflections. A diffraction vector g = 200 is indicated by an arrow in the main image of FIG. 21A. The weak-beam dark-field imaging conditions used in FIG. 21A enable the visualization of dislocation structures within the fee grains during the early stages of plastic deformation.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0216] With continued reference to FIG. 21 A, the weak-beam dark-field image reveals planar slip with precipitate shearing in the fee grains during Stage 2 of strain hardening. Dislocations in FIG. 21 A can be observed confined to specific crystallographic planes, characteristic of planar slip 106 arising from glide-plane softening. The planar slip mechanism 106 during Stage 2 can involve dislocations shearing through the Lb precipitates 104 distributed within the fee matrix 102. When dislocations shear through the Lb precipitates, the dislocations can create anti-phase boundaries within the precipitates that reduce the resistance to subsequent dislocation motion on the same slip plane. The reduction in resistance to dislocation motion on the primary slip plane can result in the glide-plane softening phenomenon, where dislocations become confined to the primary slip system rather than cross-slipping to other slip systems. The precipitate shearing mechanism observed in FIG. 21 A can contribute to the strain hardening during Stage 2 through the creation of antiphase boundaries that provide resistance to dislocation motion.

[0217] Referring to FIGS. 22A-22B, weak-beam dark-field TEM images illustrate the deformation substructure in B2 grains of the Fe-10Al-2Ti alloy during Stage 2 of the strain hardening behavior. FIG. 22A displays a weak-beam dark-field image at a scale of 200 nm, showing the microstructure of the Fe-10Al-2Ti alloy with a diffraction vector g equal to 110 indicated by an arrow. An inset (i) in FIG. 22A shows a selected area electron diffraction pattern along the

[0111] zone axis, displaying diffraction spots indexed as 110, Oil, 101, and their corresponding negative indices. The electron diffraction pattern in the inset of FIG. 22A confirms the B2 ordering of the grain, with the superlattice reflections arising from the ordered CsCl-type structure of the B2 phase.

[0218] With continued reference to FIGS. 22A-22B, FIG. 22B presents a higher magnification weak-beam dark-field image at a scale of 50 nm, also with a diffraction vector g equal to 110 indicated by an arrow. Both FIGS. 22A and 22B reveal extensive dislocation activity within the B2 phase, with dislocations appearing as bright contrast features against the darker matrix background. The extensive dislocation activity observed in FIGS. 22A and 22B demonstrates that the B2 grains can accommodate plastic strain during Stage 2 of the strain hardening behavior through dislocation motion within the ordered B2 matrix. The weak-beam dark-field images in FIGS. 22A and 22B further demonstrate that bcc precipitate shearing can occur in Stage 2 of strain hardening, where dislocations can shear through the bcc precipitates distributed within the B2 matrix.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0219] The B2 grains of the alloy compositions of the present disclosure can deform via (111) super dislocations gliding on {110} planes. The (111) super dislocations in the B2 phase can dissociate into two super partial dislocations separated by an anti-phase boundary. The dissociation of the (111) super dislocations into super partial dislocations can occur because the passage of a single (111 ) / 2 dislocation through the ordered B2 structure can create an anti-phase boundary where the ordered arrangement of atoms is disrupted. The anti-phase boundary created by the leading super partial dislocation can be eliminated by the passage of the trailing super partial dislocation, which restores the ordered B2 structure. The separation between the two super partial dislocations can be determined by the balance between the elastic repulsion of the partial dislocations and the energy of the anti-phase boundary ribbon connecting the partial dislocations.

[0220] The deformation of the B2 grains via (111) super dislocations gliding on {110} planes can enable co-deformability between the fee grains and the B2 grains during tensile loading. The {110} slip planes in the B2 phase can accommodate plastic strain in orientations that are compatible with the {111} slip planes in the fee phase, enabling strain transfer across the fcc / B2 interphase boundaries without premature damage nucleation. The co-deformability between the fee grains and the B2 grains can contribute to the high ductility achieved by the alloy compositions of the present disclosure by enabling continued plastic deformation without crack nucleation at the interphase boundaries. The extensive dislocation activity observed in the B2 phase in FIGS. 22A and 22B, combined with the planar slip and precipitate shearing observed in the fee phase in FIG. 21 A, demonstrates that both phases of the dual-phase microstructure can accommodate plastic strain during Stage 2 of the strain hardening behavior.

[0221] Referring to FIGS. 23A-23D, transmission electron microscopy (TEM) characterization illustrates the presence of B2 and bcc phases within fee grains 102 of the aged Fe-10Al-2Ti alloy. The presence of B2 and bcc phases within the fee grains 102 can represent a third microstructural constituent that distinguishes the Fe-10Al-2Ti alloy from single-precipitate alloy systems and from dual-phase alloys having precipitates only within the respective matrix grains. The B2 and bcc phases within the fee grains can contribute to Stage 4 of the strain hardening behavior through microband formation 112, as discussed with reference to FIG. 8.

[0222] FIG. 23A shows a backscattered SEM micrograph of the aged Fe-10Al-2Ti alloy displaying dark B2 grains 122 and bright fee grains 102 with the presence of dark phases 124Attorney Docket No.: MIT 26210 PCT | 88212-432499indicated by dashed box 124 within the fee grain 102. The dark phases 124 visible within the bright fee grain 102 in FIG. 23 A correspond to B2 and bee phases that have formed within the fee matrix during the aging heat treatment. The presence of these dark phases 124 within the fee grains 102 can be distinguished from the larger B2 grains that constitute the second matrix phase of the dual-phase architecture, as the dark phases within the fee grains can be smaller in size and can be distributed within the interior of the fee grains rather than forming separate grains.

[0223] With continued reference to FIGS. 23A-23D, FIG. 23B presents a selected area electron diffraction (SAED) pattern from the dark phase of the fee grain along the

[0001] zone axis. The SAED pattern in FIG. 23B shows reflections from the fundamental bcc phase and B2 ordering phase, with the B2 spots marked with arrows. The presence of both fundamental bcc reflections and B2 superlattice reflections in the SAED pattern of FIG. 23B confirms that the dark phases within the fee grains contain both disordered bcc precipitates and ordered B2 matrix regions. The B2 superlattice reflections arise from the ordered CsCl-type arrangement of atoms in the B2 structure, where the ordering creates additional diffraction spots at positions that are forbidden in the disordered bcc structure. The coexistence of bcc and B2 reflections in the SAED pattern demonstrates that the dark phases 124 within the fee grains 102 can have a similar microstructural architecture to the B2 grains discussed with reference to FIGS. 17A-17D, where disordered bcc precipitates can be distributed within an ordered B2 matrix.

[0224] FIG. 23C shows a TEM dark-field image obtained from the (200) bcc fundamental reflection with g = 200, revealing a uniform distribution of bright cuboidal shaped bcc precipitates within the dark B2 matrix. The dark-field imaging conditions used in FIG. 23C enable the visualization of the disordered bcc precipitates as bright contrast features against the darker B2 matrix background. The cuboidal morphology of the bcc precipitates shown in FIG. 23C can be consistent with the precipitate morphology observed in the B2 grains discussed with reference to FIGS. 17A-17D, where the cuboidal shape can be attributed to the elastic anisotropy of the B2 matrix and the coherent nature of the precipitate-matrix interface. The uniform distribution of bcc precipitates within the B2 regions of the fee grains shown in FIG. 23 C demonstrates that the aging heat treatment can produce a homogeneous precipitate microstructure within the B2 + bcc phases that form within the fee grains.

[0225] With continued reference to FIGS. 23A-23D, FIG. 23D displays a TEM dark-field image obtained from the (100) B2 ordering reflection with g = 100. showing dark bccAttorney Docket No.: MIT 26210 PCT | 88212-432499precipitates in the bright B2 matrix. The dark-field imaging conditions used in FIG. 23D enable the visualization of the ordered B2 matrix as bright contrast features, while the disordered bcc precipitates appear as dark contrast features because the disordered bcc structure does not contribute to the (100) superlattice reflection. The complementary contrast between FIGS. 23C and 23D, where the bcc precipitates appear bright in FIG. 23C and dark in FIG. 23D, confirms the identification of the disordered bcc precipitates within the ordered B2 matrix. The dark-field images in FIGS. 23C and 23D demonstrate that the B2 + bcc phases within the fee grains can have a hierarchical microstructure containing uniformly distributed disordered bcc precipitates within an ordered B2 matrix.

[0226] As noted above, the B2 + bcc phases within the fee grains can contribute to Stage 4 of the strain hardening behavior through the promotion of microband formation. The B2 precipitates within the fee grains can cause dislocation pileups along the primary slip system during deformation, as dislocations gliding within the fee matrix can encounter the B2 + bcc phases and can accumulate at the interfaces between the fee matrix and the B2 regions. The dislocation pileups at the B2 + bcc phases can lead to stress localization that promotes crossslip of dislocations to secondary and tertiary slip systems. The cross-slip events facilitated by the B2 + bcc phases within the fee grains can result in microband formation, where dislocations from multiple slip systems interact to form dense dislocation bands aligned along crystallographic planes. The microbands can provide long-range back stress that sustains the strain hardening rate during Stage 4, thereby delaying plastic instability and enabling the alloy to achieve high uniform elongation.

[0227] The TEM characterization results shown in FIGS. 23A-23D confirm that the aged Fe-10Al-2Ti alloy can contain B2 + bcc phases distributed within the fee grains in addition to the LL precipitates within the fee grains and the bcc precipitates within the B2 grains. The hierarchical microstructure containing multiple precipitate types within the dual-phase matrix architecture can enable the four distinct strain hardening stages discussed with reference to FIGS. 7A-7D and FIG. 8. The LL precipitates within the fee grains can govern Stage 2 through precipitate shearing and planar slip, the B2 grains containing bcc precipitates can contribute to Stage 3 through dislocation activity within the B2 matrix, and the B2 + bcc phases within the fee grains can control Stage 4 through microband formation. The sequential activation of strain hardening mechanisms associated with each microstructural constituent can enable the alloy compositions of the present disclosure to achieve both high strength and high ductility without relying on martensite or metastable austenite phases.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0228] While the Fe-10Al-2Ti was discussed above, wavy slip 106, planar slip 108, HDDWs 110, and microbands 112 can be observed for the remaining three compositions of the present disclosure. Referring to FIGS. 24A and 24B, electron channeling contrast imaging (ECCI) micrographs illustrate the deformation substructure in the Fe-6Al-6Ti alloy deformed to 13% plastic strain. FIG. 24A shows a micrograph at a scale of 1 pm depicting non-planar wavy slip deformation 106, with arrows indicating the wavy slip regions where dislocations glide on different crystallographic slip systems with a tendency to crossslip. FIG. 24B shows a higher magnification micrograph at a scale of 00 nm of the same alloy condition, further illustrating the wary slip characteristics with arrows marking the wary slip features. The strain notation in the upper right comer of each micrograph indicates the plastic strain level of 13%. The ECCI micrographs in FIGS. 24A and 24B demonstrate that the Fe-6Al-6Ti alloy exhibits predominantly non-planar wavy dislocation slip as the deformation mode during tensile loading.

[0229] The non-planar wavy slip 106 observed in the Fe-6Al-6Ti alloy can be attributed to the single-phase fee microstructure with L l2precipitates that characterizes this alloy composition. As discussed with reference to FIGS. 1 A-1E and FIG. 20, the Fe-6Al-6Ti alloy exhibits a predominantly fcc-based microstructure without the B2 phase or bcc precipitates that are present in the Fe-8Al-4Ti, Fe-10Al-2Ti, and Fe-12A1 alloy compositions. The absence of the B2 phase in the Fe-6Al-6Ti alloy can be attributed to the VEC of approximately 8.0 that stabilizes the fee phase without promoting B2 phase formation. The wavy slip mechanism in the Fe-6Al-6Ti alloy can occur when dislocations are not confined to specific {111 } planes and instead exhibit a tendency to cross-slip among multiple non-coplanar {111} planes, as illustrated schematically in FIG. 10A.

[0230] With continued reference to FIGS. 24A and 24B, the wavy slip arrangement observed in the Fe-6Al-6Ti alloy can be characteristic of alloys having medium to high stacking fault energy, where the extended dislocation width is sufficiently narrow to permit cross-slip between slip planes. The dislocations in FIGS. 24A and 24B can be distributed in a scattered pattern throughout the deforming region rather than being confined to specific crystallographic planes. The tendency of dislocations to cross-slip among multiple non-coplanar {111} planes can lead to the wavy nature of slip observed in the ECCI micrographs. The wavy slip mechanism in the Fe-6Al-6Ti alloy can contrast with the planar slip mechanism observed in the Fe-10Al-2Ti alloy during Stage 2 of strain hardening, asAttorney Docket No.: MIT 26210 PCT | 88212-432499discussed with reference to FIG. 21 A, where dislocations can be confined to the primary slip system due to glide-plane softening associated with Lh precipitate shearing.

[0231] The Fe-6Al-6Ti alloy can exhibit only two stages of strain hardening during tensile deformation, as discussed with reference to FIG. 7A. The limited number of strain hardening stages in the Fe-6Al-6Ti alloy can be attributed to the absence of the B2 grains and B2 + bcc phases within the fee grains that contribute to additional strain hardening stages in the Fe-10Al-2Ti alloy. The Fe-6Al-6Ti alloy having a single-phase fee + LL microstructure can lack the B2 grains that contribute to Stage 3 of the strain hardening behavior and can lack the B2 + bcc phases within the fee grains that control Stage 4 of the strain hardening behavior through microband formation. The absence of these microstructural constituents in the Fe-6Al-6Ti alloy can result in the limited strain hardening range observed in FIG. 6, w here the Fe-6Al-6Ti alloy exhibits a strain hardening rate that decreases from approximately 6 x 103MPa to approximately 3 x 103MPa over a true strain range of approximately 2 percent to approximately 13 percent before the curve terminates.

[0232] The ECCI observations of wavy slip in the Fe-6Al-6Ti alloy shown in FIGS. 24A and 24B can be corroborated by TEM analyses that confirm the wavy slip mechanism without deformation-induced phase transformation. TEM w eak-beam bright-field images from differently oriented

[0011] grains of the Fe-6Al-6Ti alloy deformed to 13% plastic strain can reveal the presence of dislocations without preferential confinement along specific { 111} planes. The dislocation arrangement observed in TEM, combined with the tendency of dislocations to cross-slip among multiple non-coplanar {111} planes, can lead to the wavy dislocation glide that is characteristic of wavy slip. The TEM observations of the Fe-6Al-6Ti alloy can confirm that no deformation-induced phase transformation occurs during tensile loading, as the selected area electron diffraction patterns can show only the LL superlattice reflections in addition to the fundamental fee reflections without the appearance of martensite or other transformation products.

[0233] The Fe-6Al-6Ti alloy composition can have a total composition of Al and Ti of about 12 at.%, where the Al content is approximately 6 at.% and the Ti content is approximately 6 at.%. A method of forming an alloy of the present disclosure can include providing a total composition of Al and Ti of about 12 at.%. The total composition of Al and Ti of about 12 at.% in the Fe-6Al-6Ti alloy can result in the single-phase fee + LE microstructure that exhibits wavy slip as the predominant deformation mode. The equal proportions of Al and Ti in the Fe-6Al-6Ti composition can maintain the VEC atAttorney Docket No.: MIT 26210 PCT | 88212-432499approximately 8.0, which can stabilize the fee phase without promoting B2 phase formation. The method of forming an alloy having a total composition of Al and Ti of about 12 at.% can produce different microstructures depending on the relative proportions of Al and Ti, where the Fe-6Al-6Ti composition having equal Al and Ti contents can produce a single-phase fee + LE microstructure, while the Fe-10Al-2Ti composition having unequal Al and Ti contents can produce a dual-phase fee + bcc + LL + B2 microstructure despite having the same total Al + Ti content of about 12 at.%.

[0234] The deformation behavior of the Fe-6Al-6Ti alloy characterized in FIGS. 24A and 24B can demonstrate the relationship between microstructure and strain hardening behavior in the alloy compositions of the present disclosure. The single-phase fee + LG microstructure of the Fe-6Al-6Ti alloy can provide strain hardening through dislocationprecipitate interactions during wavy slip, but can lack the additional strain hardening mechanisms provided by the B2 grains and B2 + bcc phases within the fee grains that are present in the dual-phase multi-precipitate compositions. The limited strain hardening range and reduced elongation of the Fe-6Al-6Ti alloy compared to the Fe-10Al-2Ti alloy can illustrate the advantage of the dual-phase multi-precipitate microstructure architecture for achieving both high strength and high ductility. The comparison between the Fe-6Al-6Ti alloy and the Fe-10Al-2Ti alloy can demonstrate that the presence of multiple phases and multiple precipitate types can enable additional strain hardening stages that sustain work hardening throughout the deformation process, thereby delaying plastic instability7and enabling the alloy to achieve high uniform elongation.

[0235] Referring to FIGS. 25A-25D, electron channeling contrast imaging (ECCI) micrographs and inverse pole figure (IPF) maps illustrate the deformation behavior of the Fe-8Al-4Ti alloy composition at various plastic strain levels. FIGS. 25A and 25B show ECCI micrographs of the Fe-8Al-4Ti alloy deformed to 5 percent plastic strain, corresponding to Stage 2 of the strain hardening behavior. At this strain level, the Fe-8Al-4Ti alloy can exhibit both wavy slip and planar slip deformation modes operating simultaneously within different grains. The coexistence of wavy slip and planar slip in the Fe-8Al-4Ti alloy at 5 percent strain can contrast with the predominantly planar slip observed in the Fe-10Al-2Ti alloy at the same strain level, which can be attributed to differences in the LE precipitate characteristics and phase fractions between the two compositions.

[0236] FIGS. 25C and 25D show ECCI micrographs of the Fe-8Al-4Ti alloy deformed to 17 percent plastic strain, corresponding to Stage 3 of the strain hardening behavior. AsAttorney Docket No.: MIT 26210 PCT | 88212-432499strain increases to 17 percent, high-density' dislocation walls (HDDWs) 110 can form alongside planar and wavy slip mechanisms in the Fe-8Al-4Ti alloy’. The formation of HDDWs 110 at Stage 3 in the Fe-8Al-4Ti alloy can indicate the transition from slip-dominated deformation to the development of organized dislocation structures that provide continued strain hardening. The IPF maps corresponding to the Fe-8Al-4Ti alloy at 5 percent and 17 percent plastic strain can show the distribution of deformation mechanisms across various grain orientations, demonstrating that the deformation mechanisms show no strong crystallographic dependence on grain orientation.

[0237] Referring to FIGS. 26A-26E, electron channeling contrast imaging (ECCI) micrographs and inverse pole figure (IPF) maps illustrate the deformation behavior of the Ti-free Fe-12A1 alloy composition at various plastic strain levels. FIGS. 26 A, 26B, and 26C show ECCI micrographs of the Fe-12A1 alloy deformed to 5 percent plastic strain, corresponding to Stage 2 of the strain hardening behavior. At this strain level, the Fe-12A1 alloy can exhibit wavy' slip, planar slip, and high-density dislocation walls (HDDWs) formation. The earlier appearance of HDDWs in the Fe-12A1 alloy at Stage 2 compared to the Fe-8Al-4Ti alloy at the same strain level can be attributed to differences in the precipitate characteristics and phase fractions resulting from the absence of titanium in the Fe-12A1 composition.

[0238] With continued reference to FIGS. 26A-26E, FIGS. 26D and 26E show ECCI micrographs of the Fe-12A1 alloy deformed to 22 percent plastic strain, corresponding to Stage 3 of the strain hardening behavior. At this strain level, the Fe-12A1 alloy can exhibit HDDWs and microbands (MBs) as the dominant deformation mechanisms, as shown by the respective arrows. However, the Fe-12A1 alloy does not exhibit Stage 4 deformation behavior, which can be attributed to the rapid decrease in strain hardening rate at Stage 3 that leads to early plastic instability, necking, and fracture. In contrast, the Ti-containing alloys such as Fe-6Al-6Ti, Fe-8Al-4Ti, and Fe-10Al-2Ti do not show apparent necking before fracture, which can be attributed to the sustained strain hardening behavior provided by the titanium-containing LE precipitates.

[0239] TEM analysis of the Fe-12A1 alloy deformed to 22 percent plastic strain can confirm that no deformation-induced phase transformation occurs in the fee grains, as the selected area electron diffraction (SAED) pattern shows only the LI2 superlattice and fundamental fee reflections. Weak-beam bright-field TEM images can reveal that both HDDWs and MBs are aligned along {111} planes in the Fe-12A1 alloy. The IPF mapsAttorney Docket No.: MIT 26210 PCT | 88212-432499corresponding to the Fe-12A1 alloy at 5 percent and 22 percent plastic strain can show the distribution of deformation mechanisms across various grain orientations, demonstrating that the deformation mechanisms show no strong crystallographic dependence on grain orientation. The absence of Stage 4 deformation in the Fe-12A1 alloy, combined with the rapid decrease in strain hardening rate at Stage 3, can illustrate the importance of titanium content in maintaining sustained strain hardening behavior and delaying plastic instability in the alloy compositions of the present disclosure.

[0240] Referring to FIGS. 27A-27F. microstructural characterization and mechanical property data illustrate the behavior of the recrystallized Fe-10Al-2Ti alloy without the aging treatment. The recrystallized precipitate-free Fe-10Al-2Ti alloy can provide a comparison that demonstrates the contribution of the LI 2 and B2s to the mechanical properties of the aged alloy compositions of the present disclosure. The absence of precipitates in the recrystallized alloy can result in different strain hardening behavior and reduced mechanical properties compared to the aged precipitate-containing alloy.

[0241] FIG. 27A shows a backscattered electron scanning electron microscope micrograph of the recrystallized Fe-10Al-2Ti alloy microstructure with a scale bar of 50 micrometers. The micrograph in FIG. 27A reveals a dual-phase grain structure containing fee grains and B2 grains, similar to the aged Fe-10Al-2Ti alloy discussed with reference to FIG. 15. However, the recrystallized alloy shown in FIG. 27A can lack the LI 2 precipitates within the fee grains and the bcc precipitates within the B2 grains that develop during the aging treatment. The absence of precipitates in the recrystallized alloy can be attributed to the water quenching following the recrystallization treatment, which can retain the high-temperature microstructure without allowing time for precipitate nucleation and growth.

[0242] With continued reference to FIGS. 27A-27F, FIG. 27B presents an engineering stress versus engineering strain curve for the recry stallized Fe-10Al-2Ti alloy. The vertical axis of FIG. 27B represents engineering stress measured in megapascals ranging from 0 to 1000, and the horizontal axis represents engineering strain measured in percent ranging from 0 to 40. The stress-strain curve in FIG. 27B shows that the recrystallized Fe-10Al-2Ti alloy can achieve an ultimate tensile strength approaching 1000 megapascals with elongation extending beyond 35 percent. The recrystallized Fe-10Al-2Ti alloy can achieve a yield strength of approximately 575 ± 20 MPa, an ultimate tensile strength of approximately 960 ± 25 MPa, and a total elongation of approximately 37.0 ± 2.0 percent.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0243] The mechanical properties of the recrystallized Fe-10Al-2Ti alloy can be inferior compared to the precipitate-containing aged alloy. The yield strength of approximately 575 ± 20 MPa for the recrystallized alloy can be substantially lower than the yield strength of approximately 1050 ± 50 MPa for the aged Fe-10Al-2Ti alloy. The ultimate tensile strength of approximately 960 ± 25 MPa for the recry stallized alloy can be substantially lower than the ultimate tensile strength of approximately 1650 ± 60 MPa for the aged Fe-10Al-2Ti alloy. The reduction in yield strength and ultimate tensile strength in the recrystallized alloy compared to the aged alloy can be attributed to the absence of the LG precipitates within the fee grains and the bcc precipitates within the B2 grains that provide precipitation strengthening in the aged alloy.

[0244] FIG. 27C displays a strain hardening rate versus true strain plot for the recrystallized Fe-10Al-2Ti alloy. The vertical axis of FIG. 27C represents strain hardening rate measured in megapascals ranging from 0 to 10000. and the horizontal axis represents true strain measured in percent ranging from 0 to 30. The strain hardening rate plot in FIG.27C illustrates the evolution of strain hardening behavior during deformation of the recrystallized alloy. The recrystallized Fe-10Al-2Ti alloy with no precipitates can show a continuous decrease in strain hardening rate throughout the deformation process, which can contrast with the multistage strain hardening behavior observed in the aged Fe-10Al-2Ti alloy discussed with reference to FIG. 7C.

[0245] With continued reference to FIGS. 27A-27F, the recry stallized Fe-10Al-2Ti alloy can exhibit only three stages of strain hardening during tensile deformation, which can be fewer than the four distinct strain hardening stages observed in the aged Fe-10Al-2Ti alloy. The absence of Stage 2 and Stage 4 in the recrystallized alloy can corroborate that the LG precipitates control Stage 2 and the B2 precipitates within the fee grains control Stage 4 of the strain hardening behavior. The continuous decrease in strain hardening rate observed in the recrystallized alloy can be akin to Stage 3 of the aged alloy, where the strain hardening rate decreases without the plateau regions associated with the precipitate-controlled strain hardening stages.

[0246] The increase in uniform elongation from approximately 30 percent to approximately 39 percent upon aging can indicate an enhancement of the work hardening capability of the alloy due to the introduction of Lb and B2 precipitates within the fee grains. The enhancement of work hardening capability' upon aging can contrast with the decrease in ductility observed in conventional steels upon aging. The observation that aging can increaseAttorney Docket No.: MIT 26210 PCT | 88212-432499both strength and ductility in the Fe-10Al-2Ti alloy can challenge the inverse strengthductility relationship that characterizes high-strength structural alloys. The hierarchical microstructure consisting of coherent LI 2, B2, and bcc precipitates in the fee matrix and coherent bcc precipitates in the B2 matrix can address the strength-ductility' dilemma by providing multiple strain hardening stages that sustain work hardening throughout the deformation process.

[0247] FIG. 27D shows a transmission electron microscope bright-field image with a scale bar of 2 micrometers, with arrows labeled MBs indicating the presence of microbands 112 within the deformed microstructure of the recrystallized Fe-10Al-2Ti alloy. FIG. 27E presents another transmission electron microscope image at higher magnification with a scale bar of 1 micrometer, with arrows labeled HDDWs 110 identifying high density dislocation walls formed during plastic deformation. The ECCI micrographs in FIGS. 27D and 27E reveal that microbands 112 and high density dislocation walls (HDDWs) 110 can form as prevalent deformation modes in the recrystallized Fe-10Al-2Ti alloy deformed until fracture.

[0248] With continued reference to FIGS. 27A-27F, FIG. 27F shows an inverse pole figure map plotting deformation mechanisms as a function of grain orientation at a plastic strain of 35 percent. The triangle in FIG. 27F is bounded by crystallographic orientations 001, 101, and 111, where dots represent HDDWs 110 and microbands 112. The distribution of dots across the inverse pole figure map in FIG. 27F demonstrates that both high density dislocation walls and microbands can operate across various grain orientations without strong cry stallographic dependence. The deformation mechanisms observed in the recr stallized alloy can include microbands and high densify dislocation walls as the prevalent mechanisms, with a few grains showing microband formation.

[0249] The deformation behavior of the recrystallized Fe-10Al-2Ti alloy can highlight the contribution of the B2 precipitates within the fee grams to microband formation in the aged alloy. The recrystallized alloy' lacking the B2 precipitates within the fee grains can exhibit microband formation, but the microbands in the recry stallized alloy can form without the sustained strain hardening rate associated with Stage 4 in the aged alloy. The B2 precipitates within the fee grains of the aged alloy can provide long-range back stress that sustains the strain hardening rate during Stage 4, thereby enabling the appearance of the distinct Stage 4 plateau in the strain hardening rate versus true strain plot. The absence of the Stage 4 plateau in the recrystallized alloy can demonstrate that the B2 precipitates within the fee grains canAttorney Docket No.: MIT 26210 PCT | 88212-432499be responsible for the sustained strain hardening rate during the later stages of deformation in the aged alloy.

[0250] The comparison between the recrystallized Fe-10Al-2Ti alloy discussed in FIGS.27A-27F and the aged Fe-10Al-2Ti alloy discussed above can demonstrate the contribution of the hierarchical multi-precipitate microstructure to the mechanical properties of the alloy compositions of the present disclosure. The LI 2 precipitates within the fee grains can provide precipitation strengthening that increases the yield strength from approximately 575 MPa in the recrystallized alloy to approximately 1050 MPa in the aged alloy. The LL precipitates can also control Stage 2 of the strain hardening behavior through precipitate shearing and planar slip mechanisms. The B2 precipitates within the fee grains can control Stage 4 of the strain hardening behavior through microband formation, providing long-range back stress that sustains the strain hardening rate and delays plastic instability’. The combination of Lh precipitates and B2 precipitates within the fee grains, together with the bcc precipitates within the B2 grains, can enable the aged Fe-10Al-2Ti alloy to achieve both high strength and high ductility that exceed the properties of the precipitate-free recrystallized alloy.

[0251] The alloy compositions of the present disclosure can exhibit high temperature stability characteristics that enable the alloy to be used in elevated temperature applications. The high temperature stability of the alloy compositions can be attributed to the thermal stability of the Lh precipitates within the fee grains and the bcc precipitates within the B2 grains, which can maintain the hierarchical multi-precipitate microstructure during service at elevated temperatures. The alloy compositions can be stable at a temperature approximately in a range of about 550 degrees Celsius to about 800 degrees Celsius, as discussed with reference to the thermodynamic phase diagrams and phase fraction plots.

[0252] The Lb formation temperature can be approximately 550 degrees Celsius to approximately 590 degrees Celsius, which can represent the temperature at which the LI2 precipitates begin to nucleate and grow within the fee matrix during cooling from elevated temperatures or during isothermal aging treatments. The LI2 dissolution temperature, also referred to as the solvus temperature, can be approximately 915 degrees Celsius to approximately 950 degrees Celsius, which can represent the temperature above which the LI2 precipitates dissolve into the fee matrix and the alloy transitions to a single-phase or reduced-precipitate microstructure. The LI2 precipitates can remain stable during service at temperatures below the solvus temperature, thereby maintaining the strengthening effect of the precipitates during elevated temperature applications.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0253] The alloy compositions of the present disclosure can exhibit mechanical stability at elevated temperatures, as demonstrated by Vickers hardness evolution during aging at 650 degrees Celsius. The Vickers hardness of the alloy can remain stable during extended aging times at 650 degrees Celsius, indicating that the microstructure does not undergo deleterious coarsening or phase transformation during prolonged exposure to elevated temperatures.

[0254] The alloy compositions of the present disclosure can exhibit parabolic oxidation kinetics at 600 degrees Celsius. The parabolic oxidation kinetics can indicate that the oxide layer growth rate decreases over time during oxidation at 600 degrees Celsius. The decreasing oxide layer growth rate over time can be characteristic of protective oxide formation, where the oxide layer that forms on the alloy surface can act as a diffusion barrier that slows the transport of oxygen and metal ions across the oxide layer. The parabolic oxidation kinetics can indicate that the alloy compositions of the present disclosure can form a protective oxide layer during elevated temperature exposure, which can reduce the rate of oxidation and can extend the service life of components fabricated from the alloy. The aluminum content in the alloy compositions can contribute to the formation of a protective aluminum oxide layer on the alloy surface during oxidation, where the aluminum oxide layer can provide oxidation resistance at elevated temperatures.

[0255] The alloy compositions of the present disclosure can be suitable for applications in gas turbine components. Gas turbine casings can be fabricated from the alloy compositions of the present disclosure, where the gas turbine casings can benefit from the high strength and high ductility of the alloy at elevated temperatures. Exhaust frames for gas turbines can be fabricated from the alloy compositions, where the exhaust frames can be exposed to elevated temperatures and can require both high strength and oxidation resistance. Combustor liners for gas turbines can be fabricated from the alloy compositions, where the combustor liners can be exposed to high temperatures and can require thermal stability7and oxidation resistance.

[0256] The alloy compositions of the present disclosure can be suitable for applications in steam turbine components. Steam turbine rotors can be fabricated from the alloy compositions, where the steam turbine rotors can require high strength and high ductility to withstand the mechanical stresses associated with rotation at elevated temperatures. Steam turbine blades can be fabricated from the alloy compositions, where the steam turbine blades can require high strength, high ductility, and oxidation resistance during service at elevated temperatures.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0257] The alloy compositions of the present disclosure can be suitable for applications in advanced ultra-supercritical (AUSC) power plants. AUSC power plants can operate at steam temperatures and pressures that exceed those of conventional supercritical power plants, which can require materials having high strength and oxidation resistance at elevated temperatures. Components for AUSC power plants, including turbine rotors, turbine blades, and steam piping, can be fabricated from the alloy compositions of the present disclosure. The combination of high strength, high ductility, thermal stability, and oxidation resistance exhibited by the alloy compositions can enable the alloy to be used in AUSC power plant applications where conventional steels and nickel-based superalloys can be limited by insufficient strength or excessive cost.

[0258] The alloy compositions of the present disclosure can exhibit co-deformability between the fee phase and the B2 phase during tensile loading. The co-de formability between the fee and B2 phases can enable strain partitioning between the phases without sharp strain localization at the interphase boundaries. The strain partitioning between the fee and B2 phases can occur because both phases can accommodate plastic strain through dislocation-mediated mechanisms that are compatible with one another. The fee grains can deform via dislocation glide on {111} slip planes, while the B2 grains can deform via ( 111 ) super dislocations gliding on {110} planes. The compatibility between the slip systems of the fee and B2 phases can enable strain transfer across the fcc / B2 interphase boundaries without the development of stress concentrations that could lead to premature crack nucleation.

[0259] The strain partitioning between the fee and B2 phases can result in relatively homogeneous stress accommodation throughout the dual-phase microstructure. The absence of sharp strain localization at the interphase boundaries can reduce the likelihood of damage nucleation at the fcc / B2 interfaces during tensile loading. The co-deformability' betw een the fee and B2 phases can thus contribute to the high ductility' achieved by the alloy compositions of the present disclosure by enabling continued plastic deformation without premature failure at the interphase boundaries.

[0260] The alloy compositions of the present disclosure can exhibit fracture surfaces showing cup and cone type fracture morphology'. The Fe-8Al-4Ti, Fe-10Al-2Ti, and Fe-12A1 alloy compositions can exhibit cup and cone type fracture morphology' on the fracture surfaces following tensile testing to failure. The cup and cone type fracture morphology can be characteristic of ductile fracture behavior, where the material undergoes substantial plasticAttorney Docket No.: MIT 26210 PCT | 88212-432499deformation prior to final separation. The cup and cone fracture morphology can develop through the nucleation, growth, and coalescence of microvoids during the later stages of tensile deformation, where the microvoids can nucleate at inclusions, precipitates, or other microstructural heterogeneities and can grow and coalesce to form the final fracture surface.

[0261] The ductile fracture behavior indicated by the cup and cone type fracture morphology can contrast with the intergranular fracture behavior that can occur in alloys having weak grain boundaries or brittle intermetallic phases. The Fe-6Al-6Ti alloy composition can exhibit intergranular fracture morphology on the fracture surface, which can be attributed to the different microstructure and deformation behavior of the single-phase fee + Lb composition compared to the dual-phase multi-precipitate compositions. The transition from intergranular fracture in the Fe-6Al-6Ti alloy to cup and cone type ductile fracture in the Fe-8Al-4Ti, Fe-10Al-2Ti, and Fe-12A1 alloys can be attributed to the presence of the B2 phase and the associated changes in deformation mechanisms and strain hardening behavior.

[0262] The alloy compositions of the present disclosure can exhibit excellent formability’ due to the absence of stress-assisted martensitic transformations. Stress-assisted martensitic transformations can occur in metastable austenitic alloys w here the application of mechanical stress can provide the driving force for transformation from the austenite phase to the martensite phase. The stress-assisted martensitic transformation can result in localized strain concentrations at the transformation sites, which can lead to premature damage nucleation and reduced ductility7. The alloy compositions of the present disclosure can avoid stress-assisted martensitic transformations because the fee phase can be thermodynamically stable rather than metastable, and the alloy compositions can lack the driving force for martensitic transformation during mechanical loading.

[0263] The absence of stress-assisted martensitic transformations in the alloy compositions of the present disclosure can enable the alloy to undergo plastic deformation without the formation of martensite that could act as stress concentrators or crack initiation sites. The formability7of the alloy compositions can thus be enhanced compared to metastable austenitic alloys that rely on transformation-induced plasticity (TRIP) effects for strain hardening. The alloy compositions of the present disclosure can achieve strain hardening through dislocationprecipitate interactions and the sequential activation of deformation mechanisms associated with the hierarchical multi-precipitate microstructure, rather than through deformation-induced phase transformation mechanisms.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0264] The alloy compositions of the present disclosure can exhibit excellent formability due to the absence of brittle phases within the microstructure. Brittle phases such as topologically close-packed (TCP) phases, including p and o phases, can form in alloys containing high concentrations of refractory7elements such as tungsten, molybdenum, and chromium. The TCP phases can be brittle and can act as crack initiation sites during mechanical loading, reducing the ductility and formability of the alloy. The alloy compositions of the present disclosure can avoid the formation of TCP phases by substituting vanadium and aluminum in place of tungsten, molybdenum, and chromium as bcc stabilizers, and by limiting the tantalum content to about 2 at.% to avoid TCP phase formation.

[0265] The medium to high stacking fault energy of the alloy7compositions of the present disclosure can prevent deformation twinning during tensile loading. Deformation twinning can occur in alloys having low stacking fault energy, where the extended dislocation width can be sufficiently large to permit the nucleation and propagation of twin boundaries. The twin boundaries formed during deformation twinning can act as barriers to dislocation motion and can provide strain hardening through the Hall-Petch effect associated with the reduced effective grain size. However, deformation twinning can also lead to strain localization at the twin boundaries, which can reduce the uniform elongation of the alloy.

[0266] The alloy compositions of the present disclosure can achieve strain hardening through dislocation-precipitate interactions and the formation of high-density dislocation walls and microbands, rather than through deformation twinning. The medium to high stacking fault energy' of the alloy compositions can enable dislocations to cross-slip between slip planes, which can facilitate the formation of wavy' slip, high-density dislocation walls, and microbands during tensile deformation. The cross-slip capability provided by the medium to high stacking fault energy can enable the sequential activation of deformation mechanisms that contribute to the multiple strain hardening stages, thereby enabling the alloy to achieve high uniform elongation without relying on deformation twinning for strain hardening.

[0267] The combination of co-deformability between the fee and B2 phases, the absence of stress-assisted martensitic transformations, the absence of brittle phases, and the medium to high stacking fault energy that prevents deformation twinning can contribute to the excellent formability’ of the alloy compositions of the present disclosure. The excellent formability can enable the alloy compositions to be processed using conventional metalworking techniques, including cold rolling, forging, and sheet forming, without the formation of cracks or otherAttorney Docket No.: MIT 26210 PCT | 88212-432499defects that could compromise the mechanical properties of the finished components. The ductile fracture behavior indicated by the cup and cone type fracture morphology in the Fe-8Al-4Ti, Fe-10Al-2Ti, and Fe-12A1 alloys can further demonstrate the excellent formability of the dual-phase multi-precipitate alloy compositions.

[0268] The alloy compositions of the present disclosure can achieve the combination of high strength and high ductility through the synergistic interaction of multiple microstructural elements within the hierarchical architecture. The dual-phase matrix composed of fee grains and B2 grains can provide compatible deformation pathways that enable strain accommodation throughout the polycrystalline microstructure without premature failure at interphase boundaries. The coherent precipitates distributed within the dual-phase matrix can enable sustained strain hardening through the sequential activation of deformation mechanisms at different strain levels, where each precipitate type can contribute to a distinct strain hardening stage during tensile deformation. The absence of brittle martensite phases within the microstructure can prevent premature damage nucleation that is associated with deformation-induced martensitic transformation in metastable austenitic alloys.

[0269] The dual-phase matrix architecture can provide compatible deformation pathways through the co-deformability of the fee and B2 phases. The fee grains can deform via dislocation glide on {111} slip planes, while the B2 grains can deform via (111) super dislocations gliding on { 110} planes. The slip systems of the fee and B2 phases can be geometrically compatible, enabling strain transfer across the fcc / B2 interphase boundaries without the development of stress concentrations that could lead to crack nucleation. The similar hardness values of the fee and B2 phases can enable both phases to carry comparable loads during deformation, resulting in relatively homogeneous stress accommodation throughout the dual-phase microstructure. The homogeneous stress accommodation can reduce the likelihood of strain localization at the interphase boundaries, thereby enabling continued plastic deformation without premature failure.

[0270] The coherent LL precipitates within the fee grains can contribute to strain hardening during the early stages of plastic deformation. The Lh precipitates can interact with dislocations through shearing mechanisms, where dislocations can cut through the precipitates and create anti-phase boundaries within the ordered LI 2 structure. The creation of anti-phase boundaries can provide resistance to dislocation motion and can contribute to strain hardening. The shearing of LL precipitates by dislocations can also result in glideplane softening, where the destruction of the ordered structure on the primary’ slip plane canAttorney Docket No.: MIT 26210 PCT | 88212-432499reduce the resistance to subsequent dislocation motion on the same plane. The glide-plane softening can confine dislocations to the primary slip system, resulting in planar slip that characterizes the early stages of strain hardening in the alloy compositions of the present disclosure.

[0271] The coherent bcc precipitates within the B2 grains can contribute to strain hardening during intermediate stages of plastic deformation. The bcc precipitates can interact with dislocations gliding within the B2 matrix, providing obstacles to dislocation motion that contribute to strain hardening. The shearing of bcc precipitates by super dislocations within the B2 matrix can provide additional resistance to dislocation motion. The co-deformability between the fee grains containing Lk precipitates and the B2 grains containing bcc precipitates can enable continued strain hardening as the deformation transitions from the LI 2-con trolled regime to the B2-controlled regime.

[0272] The B2 precipitates within the fee grains can contribute to strain hardening during the later stages of plastic deformation through the promotion of microband formation. The B2 precipitates within the fee grains can cause dislocation pileups along the primary slip system, where dislocations gliding within the fee matrix can accumulate at the interfaces between the fee matrix and the B2 precipitates. The dislocation pileups can generate long-range back stress that can activate secondary’ and tertiary slip systems through cross-slip mechanisms. The activation of multiple slip systems can result in the formation of microbands, where dislocations from multiple slip systems interact to form dense dislocation bands aligned along crystallographic planes. The microbands can provide sustained strain hardening during the later stages of deformation by maintaining the back stress that delays plastic instability.

[0273] The sequential activation of deformation mechanisms associated with the LL precipitates, the B2 grains, and the B2 precipitates within the fee grains can enable the alloy compositions to exhibit multiple distinct strain hardening stages during tensile deformation. The transition between strain hardening stages can occur as the strain hardening contribution from one microstructural element begins to saturate and the contribution from another microstructural element becomes dominant. The multiple strain hardening stages can sustain work hardening throughout the deformation process, delaying the onset of necking instability and enabling the alloy to achieve high uniform elongation before failure.Attorney Docket No.: MIT 26210 PCT | 88212-432499

[0274] The compositional control through Al and Ti content can enable tuning of phase fractions and mechanical properties in the alloy compositions of the present disclosure. Increasing the Al content while decreasing the Ti content can promote the formation of the B2 phase at the expense of the fee phase, as the increased Al content can lower the VEC and can stabilize the fee + bcc phase field. The variation in B2 phase fraction with Al and Ti content can influence the strain hardening behavior and mechanical properties of the alloy compositions. Alloy compositions having low B2 phase fractions can exhibit fewer strain hardening stages and reduced elongation, while alloy compositions having intermediate B2 phase fractions can exhibit multiple strain hardening stages and high elongation. Alloy compositions having high B2 phase fractions can exhibit reduced strengthductility combinations compared to compositions having intermediate B2 phase fractions.

[0275] The Ti content can influence the anti-phase boundary energy of the Lk precipitates, which can affect the shearability of the precipitates and the deformation mechanisms operating during the early stages of strain hardening. Higher Ti content can increase the antiphase boundary' energy' of the LL precipitates, which can reduce the shearability' of the precipitates and can inhibit planar slip. The inhibition of planar slip at higher Ti content can lead to the formation of dislocation walls and microbands at earlier strain levels, which can degrade the mechanical properties by reducing the strain hardening range. The Ti content can thus be selected to optimize the balance between LL precipitate stability' and precipitate shearability', enabling the alloy to achieve the multiple strain hardening stages that provide the combination of high strength and high ductility.

[0276] The total composition of Al and Ti can be maintained within a range that enables the formation of the hierarchical multi-precipitate microstructure. The total Al + Ti content can influence the volume fraction of LL precipitates within the fee grains and the volume fraction of bcc precipitates within the B2 grains. The total Al + Ti content can be selected to provide sufficient precipitate volume fractions for effective strengthening while avoiding excessive precipitate coarsening or the formation of deleterious phases. The compositional control through Al and Ti content can thus enable the alloy compositions of the present disclosure to be tuned for different mechanical property' requirements, where the relative proportions of Al and Ti can be adjusted to achieve different balances between strength and ductility.

[0277] Examples of the above-described embodiments can include the following:Attorney Docket No.: MIT 26210 PCT | 88212-4324991. A composition, comprising:an alloy having:iron (Fe) approximately in a range of about 35 at.% to about 40 at.%; cobalt (Co) approximately in a range of about 12 at.% to about 17 at.%; nickel (Ni) approximately in a range of about 28 at.% to about 32 at.%; aluminum (Al) approximately in a range of about 5 at.% to about 14 at.%; tantalum (Ta) approximately in a range of about 1 at.% to about 2.5 at.%; titanium (Ti) approximately in a range of about 0 at.% to about 7 at.%; vanadium (V) approximately in a range of about 3 at.% to about 8 at.%; and boron (B) approximately in a range of about 0.01 at.% to about 0.3 at.%, wherein a total composition of Al and Ti is approximately in a range of about 5 at.% to about 14 at%.2. The composition of example 1, wherein the alloy composition is stable at a temperature approximately in a range of about 550 degrees Celsius to about 800 degrees Celsius.3. The composition of example 1 or example 2, wherein the Ni is approximately in a range of about 29.98 at.% to about 31 at.%.4. The composition of any of examples 1 to 3, wherein the Co is approximately in a range of about 14 at.% to about 15 at.%.5. The composition of any of examples 1 to 4, wherein the Al is approximately in a range of about 7 at.% to about 10 at.%.6. The composition of any of examples 1 to 5, wherein the Ta is at about 2 at.%.7. The composition of any of examples 1 to 6, wherein the Ti is at about 2 at.%.8. The composition of any of examples 1 to 7, wherein the V is at about 5 at.%.9. The composition of any of examples 1 to 8, wherein the B is at about 0.02 at.%.10. The composition of any of examples 1 to 9, wherein the B is included as a master alloy of Fe-B to prevent evaporation during casting.Attorney Docket No.: MIT 26210 PCT | 88212-43249911. A method of forming an alloy, comprising:combining an iron-aluminum-titanium (Fe-Al-Ti) alloy system with:a cobalt (Co) approximately in a range of about 12 at.% to about 17 at. %; a nickel (Ni) approximately in a range of about 28 at.% to about 32 at.%; a tantalum (Ta) approximately in a range of about 1 at.% to about 2.5 at.%; a vanadium (V) approximately in a range of about 3 at.% to about 8 at.%; and a boron (B) approximately in a range of about 0.01 at.% to about 0.3 at.%, to form an alloy system,wherein a total composition of Al and Ti is approximately in a range of about 5 at.% to about 14 at%.12. The method of example 11, further comprising adjusting an amount of one or more of the Co, Ni, Ta, V, or B to change phase stability of the alloy system.13. The method of example 11 or example 12, further comprising adjusting an amount of one or more of Al or Ti in the alloy system.14. The method of example 13, wherein adjusting further comprises increasing the amount of Al or decreasing the amount of Ti.15. The method of any of examples 11 to 14, wherein the Ni is approximately in a range of about 29.98 at.% to about 31 at.%.16. The method of any of examples 11 to 15, wherein the Co is approximately in a range of about 14 at.% to about 15 at.%.17. The method of any of examples 11 to 16, wherein the B is added as a master alloy of Fe-B to prevent evaporation during casting.18. The method of any of examples 11 to 17, wherein the Al is approximately in a range of about 7 at.% to about 10 at.%.19. The method of any of examples 11 to 18, wherein the Ta is at about 2 at. %.Attorney Docket No.: MIT 26210 PCT | 88212-43249920. The method of any of examples 11 to 19, wherein the Ti is at about 2 at.%.21. The method of any of examples 11 to 20, wherein the V is at about 5 at.%.22. The method of any of examples 11 to 21, wherein the B is at about 0.02 at.%.23. The method of any of examples 11 to 22, wherein the total composition of Al and Ti is about 12 at%.24. A method of use of the composition of examples 11 to 23, wherein the alloy system is stable at a temperature approximately in a range of about 550 degrees Celsius to about 800 degrees Celsius.25. A composition, comprising:an alloy that includes one or more of a face-centered cubic phase that contains LI 2 precipitates distributed therein, or a B2 phase that contains body-centered cubic precipitates distributed therein.26. The composition of example 25, wherein the alloy comprises a dual-phase matrix architecture composed of the face-centered cubic phase and the B2 phase.27. The composition of example 25 or example 26, wherein the face-centered cubic phase further contains B2 precipitates distributed therein.28. The composition of any of examples 25 to 27, wherein the alloy is free from metastable austenite phases and martensite phases.29. The composition of any of examples 25 to 28, wherein the LL precipitates are coherent with the face-centered cubic phase.30. The composition of example 29, wherein a lattice misfit between the LL precipitates and the face-centered cubic phase is approximately -0.84%.31. The composition of any of examples 25 to 30. wherein the body-centered cubic precipitates are coherent wi th the B2 phase.Attorney Docket No.: MIT 26210 PCT | 88212-43249932. The composition of example 31, wherein a lattice misfit between the body-centered cubic precipitates and the B2 phase is approximately -0.26%.33. The composition of any of examples 25 to 32, wherein the body-centered cubic precipitates within the B2 phase have a cuboidal morphology.34. The composition of any of examples 25 to 33, wherein the alloy exhibits at least three distinct strain hardening stages during tensile deformation.35. The composition of any of examples 25 to 34, wherein the alloy is formed from a master alloy of iron and boron combined with additional alloying elements.36. The composition of example 35, wherein the additional alloying elements include nickel, cobalt, aluminum, titanium, tantalum, and vanadium.37. A method of forming an alloy, comprising:combining iron, nickel, cobalt, aluminum, and vanadium to form an alloy system; subjecting the alloy system to a recrystallization treatment; andsubjecting the alloy system to an aging treatment to precipitate Lfi precipitates within face-centered cubic grains and body -centered cubic precipitates within B2 grains, thereby forming a dual-phase matrix architecture.38. The method of example 37, wherein combining iron, nickel, cobalt, aluminum, and vanadium further comprises combining with titanium, tantalum, and boron.39. The method of example 38, wherein the boron is included as a master alloy of ironboron to prevent evaporation during casting.40. The method of any of examples 37 to 39, wherein the recrystallization treatment is performed at a temperature in a range of about 1150 degrees Celsius to about 1190 degrees Celsius.41. The method of example 40, wherein the recrystallization treatment is performed for a duration in a range of about 15 seconds to about 20 minutes.Attorney Docket No.: MIT 26210 PCT | 88212-43249942. The method of any of examples 37 to 41, wherein the aging treatment is performed at a temperature of approximately 650 degrees Celsius.43. The method of example 42, wherein the aging treatment is performed for a duration in a range of about 10 hours to about 30 hours.44. The method of any of examples 37 to 43, further comprising cold rolling the alloy system with approximately 60 percent reduction in thickness prior to the recrystallization treatment.45. The method of any of examples 37 to 44, further comprising homogenizing the alloy system at approximately 1150 degrees Celsius for approximately 10 hours prior to cold rolling.46. The method of example 45, wherein the alloy further comprises iron in a range of about 35 at.% to about 40 at.%, nickel in a range of about 28 at.% to about 32 at.%, cobalt in a range of about 12 at.% to about 17 at.%, aluminum in a range of about 5 at.% to about 14 at.%, and vanadium in a range of about 3 at.% to about 8 at.%.47. The method of example 46, wherein the alloy further comprises titanium in a range of about 0 at.% to about 7 at.% and tantalum in a range of about 1 at.% to about 2.5 at.%.

[0278] Although the procedures provided for herein are described in conjunction with performing thermomechanical processing of Fe-Al-Ti alloy systems to achieve dual -phase multi-precipitate microstructures, the instruments and procedures provided for herein can also be used and applied in other alloy design and processing applications, including the development of high-strength structural components for automotive, aerospace, energy' generation, and defense applications. A person skilled in the art, in view' of the present disclosures, will understand how' other alloy design and processing applications, including the development of high-strength structural components for automotive, aerospace, energy generation, and defense applications can be implemented in view of the present disclosures.

[0279] One skilled in the art will appreciate further features and advantages of the disclosure based on the above-described embodiments. Accordingly, the disclosure is not to be limited by what has been particularly show n and described, except as indicated by the appended claims. By way of example, the alloy compositions and thermomechanicalAttorney Docket No.: MIT 26210 PCT | 88212-432499processing methods disclosed herein can be adapted for use in fabricating components for nuclear reactor applications, cryogenic storage vessels, or high-temperature fasteners where combinations of high strength and high ductility are required. A person skilled in the art, in view of the present disclosures, will be able to adapt some or all of the various systems, devices, and methods disclosed herein for developing precipitation-strengthened alloys in other base metal systems, optimizing heat treatment schedules for different component geometries, or tailoring microstructural features for specific service environments and loading conditions. All publications and references cited herein are expressly incorporated herein by reference in their entirety.

Claims

1. Attorney Docket No.: MIT 26210 PCT | 88212-432499CLAIMSWhat is claimed is:

1. A composition, comprising:an alloy having:iron (Fe) approximately in a range of about 35 at.% to about 40 at.%; cobalt (Co) approximately in a range of about 12 at.% to about 17 at.%; nickel (Ni) approximately in a range of about 28 at.% to about 32 at.%; aluminum (Al) approximately in a range of about 5 at.% to about 14 at.%; tantalum (Ta) approximately in a range of about 1 at.% to about 2.5 at.%; titanium (Ti) approximately in a range of about 0 at.% to about 7 at.%; vanadium (V) approximately in a range of about 3 at.% to about 8 at.%; and boron (B) approximately in a range of about 0.01 at.% to about 0.3 at.%, wherein a total composition of Al and Ti is approximately in a range of about 5 at.% to about 14 at%.

2. The composition of claim 1, wherein the alloy composition is stable at a temperature approximately in a range of about 550 degrees Celsius to about 800 degrees Celsius.

3. The composition of claim 1 , wherein the Ta is at about 2 at.%.

4. The composition of claim 1, wherein the Ti is at about 2 at.%.

5. The composition of claim 1, wherein the V is at about 5 at.%.

6. The composition of claim 1, wherein the B is at about 0.02 at.%.

7. The composition of claim 1, wherein the B is included as a master alloy of Fe-B to prevent evaporation during casting.

8. A method of forming an alloy, comprising:combining an iron-aluminum-titanium (Fe-Al-Ti) alloy system with:a cobalt (Co) approximately in a range of about 12 at.% to about 17 at.%; a nickel (Ni) approximately in a range of about 28 at.% to about 32 at.%; a tantalum (Ta) approximately in a range of about 1 at.% to about 2.5 at.%;Attorney Docket No.: MIT 26210 PCT | 88212-432499a vanadium (V) approximately in a range of about 3 at.% to about 8 at.%; and a boron (B) approximately in a range of about 0.01 at.% to about 0.3 at.%, to form an alloy system,wherein a total composition of Al and Ti is approximately in a range of about 5 at.% to about 14 at%.

9. The method of claim 8, further comprising adjusting an amount of one or more of the Co, Ni, Ta, V, or B to change phase stability of the alloy system.

10. The method of claim 8, further comprising adj usting an amount of one or more of Al or Ti in the alloy system.

11. The method of claim 10, wherein adjusting further comprises increasing the amount of Al or decreasing the amount of Ti.

12. The method of claim 8, wherein the B is added as a master alloy of Fe-B to prevent evaporation during casting.

13. The method of claim 8, wherein the total composition of Al and Ti is about 12 at%.

14. A composition, comprising:an alloy that includes one or more of a face-centered cubic phase that contains Lh precipitates distributed therein, or a B2 phase that contains body-centered cubic precipitates distributed therein.

15. The composition of claim 14, wherein the alloy comprises a dual-phase matrix architecture composed of the face-centered cubic phase and the B2 phase.

16. The composition of claim 14, wherein the face-centered cubic phase further contains B2 precipitates distributed therein.

17. The composition of claim 14, wherein the Lb precipitates are coherent with the facecentered cubic phase.Attorney Docket No.: MIT 26210 PCT | 88212-43249918. The composition of claim 14, wherein the body-centered cubic precipitates are coherent with the B2 phase.

19. The composition of claim 14, wherein the body-centered cubic precipitates within the B2 phase have a cuboidal morphology .

20. The composition of claim 14, wherein the alloy exhibits at least three distinct strain hardening stages during tensile deformation.

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