A method for alloy surface hardening by phase tuning via interstitial uptake

Abstract

Interstitial uptake of a metal alloy is leveraged to either stabilize low-temperature miscibility gaps or promote new miscibility gaps with respect to the interstitial elements at temperatures above the native metal interstitial-free alloy miscibility gap. The miscibility gap pathway offers selective precipitation at the surface while maintaining base alloy properties throughout the bulk of an article or component. Interstitially-induced separated phases offer even higher control over the interdiffusion between the metal alloy and a coating, which will expand the number of compatible coatings for a given alloy, increasing the potential performance of the alloy in harsh environments.

Description

Ref. No.: JHU-42257.601A METHOD FOR ALLOY SURFACE HARDENING BY PHASE TUNING VIA INTERSTITIAL UPTAKECROSS-REFERENCE TO RELATED APPLICATIONS

[0001] This application claims the benefit of U.S. Provisional Application No. 63 / 574,943, filed April 5, 2024, the contents of which are incorporated herein by reference in its entirety.FEDERALLY SPONSORED RESEARCH OR DEVELOPMENT

[0002] This invention was made with government support under grant N00014-20- 1-2368 awarded by the Office of Naval Research. The government has certain rights in the invention.BACKGROUND

[0003] Refractory multi-principal element alloys (RMPEAs), also called high entropy alloys (HEAs) or complex concentrated alloys (CCAs), consist of a combination of high melting point elements (Hf, Mo, Nb, Ta, Ti, V, W, Rh, and Zr) in near equimolar proportions, and promise strength and stability at high temperatures beyond that of incumbent dilute alloys (Ni and Co- based superalloys, Nb-based alloys). Liu et al., 2022; Senkov et al., 2018a; Senkov et al., 2019; Senkov et al, 2018b; An et al., 2021; Zyka et al., 2019.

[0004] While the major motivation for developing this class of alloys is to enable hotter jet engines with cleaner emissions and higher efficiency, there has been increased interest in these alloys as metallic biomaterials. Todai et al., 2017; Whitfield et al., 2021. To meet the needs of target applications of jet engines, gas turbines for electricity generation and naval propulsion, hypersonic flight, and orthopedic implants, RMPEAs must simultaneously satisfy high-temperature creep, room-temperature ductility, toughness, corrosion, and oxidation standards.

[0005] Understanding and predicting the phase evolution during oxidation in RMPEAs, especially those with Group IV and V elements, has proven difficult due to the ambiguity of how elements with vastly different oxygen solubilities (i.e., Nb: 2.9 at% O; a-Ti: 33 at% O; a-Zr: 30 at% O at 1050 °C) (Naito and Matsui, 1984; Murray and Wriedt, 1987; Abriata etRef. No.: JHU-42257.601 al., 1986; Gebhardt and Rothenbacher, 1963), oxygen affinities, and cation diffusivities behave when combined in a non-dilutc, heterogeneous chemical landscape. An understanding of the phase evolution phenomenon in these alloys under oxidation would also provide insight into other alloy compositions and uptake of other interstitial anions with similar thermodynamic and kinetic disparities amongst the constituent elements.SUMMARY

[0006] Certain aspects of the presently disclosed subject matter having been stated hereinabove, which are addressed in whole or in part by the presently disclosed subject matter, other aspects will become evident as the description proceeds when taken in connection with the accompanying Examples and Figures as best described herein below.

[0007] In some aspects, the present disclosure relates to a method of preparing an interstitially-induced miscibility gap resulting in separation of phases on or near a surface of a metal alloy, said method comprising: heating a metal alloy to an elevated temperature that may introduce ordering or clustering of metallic elements of said metal alloy; and introducing at least one interstitial element at the elevated temperature, which is promoted or excluded in certain interstitial lattice sites due to the metallic ordering or clustering of the metallic elements of said metal alloy, wherein two or more different phases are separated due to interstitially-induced immiscibility, and wherein increased interstitial uptake at elevated temperature may promote a sequence of further phase separations of the initially- separated phases, resulting in a hierarchy of multiple phase separations as interstitial uptake proceeds.

[0008] In some other aspects, a metal alloy comprising interstitially-induced separated phases on or near a surface of a metal alloy is described.BRIEF DESCRIPTION OF THE FIGURES

[0009] The patent or application file contains at least one drawing executed in color. Copies of this patent or patent application publication with color drawing(s) will be provided by the Office upon request and payment of the necessary fee.Ref. No.: JHU-42257.601

[0010] Having thus described the presently disclosed subject matter in general terms, reference will now be made to the accompanying Figures, which arc not necessarily drawn to scale, and wherein:

[0011] FIG. 1 illustrates a temperature-composition phase diagram with a simple case of the phenomenon of alloy miscibility gaps, which are regions that result in separation into two or more distinct compositions, local atomic bonding structures, and / or crystal structures.

[0012] FIG. 2 is a schematic of an embodiment of surface spinodal decomposition during heat treatment of equiatomic NbTiZr in the presence of oxygen, wherein at least two different phases are produced, which stabilized the precipitation of hard oxide compounds which further enhance hardness and surface wear-resistance.

[0013] FIG. 3A shows a scanning electron microscopy (SEM)-backscattered electrons (BSE) image of HfNbTiZr after 30 min of oxidation showing a similar hierarchical microstructure to NbTiZr. The inset shows a higher magnification of discontinuous “banding” of the Ti-rich phase rather than the continuous scale seen in NbTiZr.

[0014] FIG. 3B shows a SEM-BSE of the oxygen diffusion zone (ODZ) adjacent to the base metal, showing phase separation and a similar darker, more oxygen-rich region as in NbTiZr.

[0015] FIG. 3Cis a higher magnification colorized (yellow is highest intensity and black is lowest) SEM-BSE of the interface between layers 1 and 2, which reveals the same hierarchical phase separation as NbTiZr but with a micro structure more closely resembling nucleation and growth.

[0016] FIG. 3D shows that the average hardness and modulus plotted for each x-distance reveals a similar exceptional subsurface hardening phenomenon resulting from the hierarchical phase separation. The shaded area displayed is the standard deviation.

[0017] FIG. 3E shows contour plots for the modulus and hardness of the indentation region, with a highest hardness value of 21.4 GPa in the Ti-rich banded region.

[0018] FIG. 4A provides the indentation depths for each hardness / modulus contour plot in FIG. 7. The grayscale contour plots show that the indentation depths sufficiently exceed the limit for indentation size effects from the Berkovich tip, ranging from 415-820 nm in depth. This indentation image is from the 10-minute NbTiZr microstructure.

[0019] FIG. 4B is analogous to FIG. 4A but is from the 30-minute NbTiZr micro structure.Ref. No.: JHU-42257.601

[0020] FIG. 4C is analogous to FIG. 4A but is from the 3-hour NbTiZr microstructure.

[0021] FIG. 4D is analogous to FIG. 4A but is from the 30-minutc HfNbTiZr microstructure.

[0022] FIG. 5A shows the row-averaged hardness and modulus, with the surface lying at the right of the x-axis, of equimolar NbTiZr oxidized at 1250°C for 30 minutes and then crosssectioned.

[0023] FIG. 5B shows the row-averaged hardness and modulus, with the surface lying at the right of the x-axis, of equimolar HfNbTaTiZr oxidized at 1250°C for 30 minutes and then cross- sectioned.

[0024] FIG. 5C shows the row-averaged hardness and modulus, with the surface lying at the right of the x-axis, of 32.3Nb-32.3Ti-32.3Zr-3Al (at.%) oxidized at 1250°C for 30 minutes and then cross-sectioned.

[0025] FIG. 5D shows the row-averaged hardness and modulus, with the surface lying at the right of the x-axis, of 39Nb-39Ti-19Zr-3Al (at.%) oxidized at 1250°C for 30 minutes and then cross-sectioned.

[0026] FIG. 5E shows the row-averaged hardness and modulus, with the surface lying at the right of the x-axis, of 39Nb-39Ti-9.5Zr-9.5Hf-3Al (at.%) oxidized at 1050°C for 1 hour and then cross-sectioned.

[0027] FIG. 5F shows the row-averaged hardness and modulus, with the surface lying at the right of the x-axis, of 39Nb-39Ti-9.5Zr-9.5Hf-3Al (at.%) oxidized at 1250°C for 30 minutes and then cross- sectioned.

[0028] FIG. 6A is the average modulus versus x-distance for 5 pm and 10 qm indent spacing from the layered surface micro structure of NbTiZr after 10-minutes of oxidation at 1050°C, showing no significant variation in the measurements. This confirms that the modulus measurements are not inflated by indent interactions.

[0029] FIG. 6B is a SEM-SE image of an indent at the interface of the base metal and spinodal region from the 10-minute micro structure (left). Shear banding (yellow a rows) is visible in the base metal portion only, both on the cross-sectional sample surface and within the indent. Homogeneous deformation of the spinodal phases is clearly shown. An SEM-SE image of the uncoarsened three-phase region showing homogeneous deformation and no cracking is shown on the right.Ref. No.: JHU-42257.601

[0030] FIG. 6C shows the average hardness versus x-distance for 5 pm and 10 pm indent spacing from the 10-minutc microstructure, showing no significant variation in the measurements. This confirms that the hardness measurements are not inflated by indent interactions.

[0031] FIG. 6D is a SEM-SE image of an indent in the base metal from the 10-minute microstructure, with significant shear banding (yellow arrows) visible on the surface and within the indent. (Right) SEM-SE image of the outer oxides that shows no cracking and homogeneous deformation despite high hardness and a heterogeneous micro structure of phases typically expected to be brittle.

[0032] FIG. 7A is a pair distribution function (PDF) analysis of the change in short-range order between 1250 °C and 1050 °C, with the gray line denoting the difference in the curves and X=Nb,Ti.

[0033] FIG. 7B is a SEM-BSE image showing the microstructure after oxidation at 1050 °C with single phase base metal on the left and spinodal decomposition on the right. Insets from left to right show scanning transmission electron microscopy (STEM)-energy dispersive X- ray spectroscopy (EDS) of the spinodal decomposition, and SEM-BSE of spinodal decomposition consuming a pocket of base metal.

[0034] FIG. 7C is a TEM Selected Area Diffraction of the base metal and spinodal decomposition. Yellow arrows indicate arcs of diffuse intensity indicative of SRO. Magnified images show peak splitting indicative of spinodal decomposition.

[0035] FIG. 7D is a STEM-EDS line scan of the spinodal decomposition as shown by the arrow in (B), with an integrated width of 16nm.

[0036] FIG. 7E is a SEM-BSE image showing the onset of coarsening of the spinodal into lamellae at the base metal interface after 30 minutes of oxidation. The inset shows a higher magnification of the coarsening reaction.

[0037] FIG. 7F is a SEM-BSE image showing that the innermost layer adjacent to the base metal is fully coarsened after 3 hours of oxidation.

[0038] FIG. 8A shows the evolution of mechanical properties with microstructural variations across oxidation times. Nanoindentation plots of average elastic modulus and hardness for each x-distance, which is clearly portrayed in the corresponding two- dimensional heat maps. The standard deviation for each x-distance is shown as a shadedRef. No.: JHU-42257.601 error region in the plots. Nanoindentation was performed on representative microstructures observed after 10 min oxidation at 1050 °C, collected from an identical uncoarscncd region in the 30 min sample. FIG. 8A extends almost exactly to the surface.

[0039] FIG. 8B is analogous to FIG. 8A but nanoindentation was performed on representative microstructures observed after 30 min oxidation at 1050 °C, collected from a coarsened region. FIG. 8B is limited to a subsurface range focusing on the coarsening front.

[0040] FIG. 8C is analogous to FIG. 8A but nanoindentation was performed on representative microstructures observed after 3 hours oxidation at 1050 °C. FIG. 8C extends almost exactly to the surface.

[0041] FIG. 9A shows the driving force for OJ transformation and resultant coarsening. FIG. 9A is a SEM-BSE with adjusted gamma showing the coarsening front in the 30-min sample and a representative area of phases captured in FIG. 9B. The emergence of a (Ti,O)-rich scale at this front is evident.

[0042] FIG. 9B is a STEM-HAADF of the 2-phase lamellar interface from a location in the 3-hour sample equal distance from the base metal as the location shown in FIG. 9A, with corresponding fast Fourier transforms of the Nb-rich body-centered cubic (BCC) and Zr-rich co phases.

[0043] FIG. 9C is a geometric phase analysis (GPA) eXx map of the BCC-co interface showing highly localized strain fields at the misfit dislocations.

[0044] FIG. 9D is a colorized inverse fast Fourier transform of a larger view field of FIG. 9B that shows misfit dislocation spacing. Masked reflections are (211)[3 and (l lOl)co.

[0045] FIG. 10A shows the hierarchical structure after 3 hours of oxidation displayed across length scales. A SEM-BSE of microstructure extending from the base metal to the surface. Region II exhibits a 2-phase lamellar layer with the Nb-rich A2 phase and a (Zr,O)-rich co phase (as shown in FIG. 9), with Ti relatively even between them. Region III shows a 3- phase lamellar structure due to a secondary phase separation within the co lamella of Region II into approximately 10 nm sublamellae with Zr / Ti segregation (Transformation 3). Region IV contains a finely-dispersed microstructure with oxides.

[0046] FIG. 10B left to right: STEM-electron energy loss spectroscopy (EELS) of Region II (color legend for all chemical mapping displayed); STEM-EDS of Region III with STEM- Fast Fourier Transform (FFT) of Nb-rich BCC and representative location; and STEM-Ref. No.: JHU-42257.601HAADF of the Regions III-IV interface (coarsening front, marked by white line) with inset FFT of the oj-TixOyinternal scale and inset of STEM-EDS map of Region IV.

[0047] FIG. 10C is a colorized STEM-high-angle annular dark-field imaging (HAADF) (adjusted gamma, radial difference filter applied) from Region III showing nanotwins in the Zr sublamella with Zr-rich FFT from location shown and Ti-co suboxide FFT from a representative location, STEM-EDS line scan of the sublamellae in Region III as shown by the arrow in FIG. 10B with an integrated width of 27 nm.

[0048] FIG. 11 A is a TEM micrograph of the layer adjacent to the base metal in a Nb39- Ti39-Zrl9-3A1 (at%) alloy after oxidation at 1050°C for 1 hour. It reveals a morphology of globular phases within a continuous matrix.

[0049] FIG. 1 IB is a TEM selected area diffraction of the phases in FIG. 11 A. It shows a disordered BCC phase (blue) with two crystallographic variants of the co phase (green and orange). Additional selected area diffraction with a smaller aperture confirmed that the co reflections of both valiants are from the globular phases, while the BCC reflections arc from the matrix.

[0050] FIG. 11C shows STEM-EDS maps of the distribution of Nb, Ti, Zr, and O from another location in the same microstructure as shown in FIG. 11A and FIG. 1 IB. It can be seen that Nb segregates to the matrix while Zr, Ti, and O segregate to the co phase as in equimolar NbTiZr.

[0051] FIG. 12 illustrate scanning electron microscope backscattered electron images of the 80% cold-rolled and oxidized NbTiZr. The thinned samples targeted the center microstructure.

[0052] FIG. 13 shows load-displacement curves from eight Berkovich tip nanoindentations in the center of the 5mm cube of NbTiZr oxidized at 1050°C for 30 minutes. This corresponds to the “Treated NbTiZr - Center” hardness and modulus values in Table 1.

[0053] FIG. 14 displays scanning electron microscope images collected by secondary electron detectors. It demonstrates that the molybdenum-containing alloy mentioned herein forms porosity in the surface layers when the disclosed method is applied in the presence of volatile phase formers such as Mo. The surface is oriented at the top of the image, and pores are easily observed as the black regions within the sample.Ref. No.: JHU-42257.601

[0054] FIG. 15 shows bar graphs summarizing the average and standard error (error bars) of the livc / dcad fluorescence intensity. The results arc normalized to the average intensity after one day of culturing for each group.DETAILED DESCRIPTION

[0055] The presently disclosed subject matter now will be described more fully hereinafter with reference to the accompanying Figures, in which some, but not all embodiments of the inventions are shown. Like numbers refer to like elements throughout. The presently disclosed subject matter may be embodied in many different forms and should not be construed as limited to the embodiments set forth herein; rather, these embodiments are provided so that this disclosure will satisfy applicable legal requirements. Indeed, many modifications and other embodiments of the presently disclosed subject matter set forth herein will come to mind to one skilled in the art to which the presently disclosed subject matter pertains having the benefit of the teachings presented in the foregoing descriptions and the associated Figures. Therefore, it is to be understood that the presently disclosed subject matter is not to be limited to the specific embodiments disclosed and that modifications and other embodiments are intended to be included within the scope of the appended claims.

[0056] Following long-standing patent law convention, the terms “a,” “an,” and “the” refer to “one or more” when used in this application, including the claims. Thus, for example, reference to “a subject” includes a plurality of subjects, unless the context clearly is to the contrary (e.g., a plurality of subjects), and so forth.

[0057] Throughout this specification and the claims, the terms “comprise,” “comprises,” and “comprising” are used in a non-exclusive sense, except where the context requires otherwise. Likewise, the term “include” and its grammatical variants are intended to be nonlimiting, such that recitation of items in a list is not to the exclusion of other like items that can be substituted or added to the listed items.

[0058] For the purposes of this specification and appended claims, unless otherwise indicated, all numbers expressing amounts, sizes, dimensions, proportions, shapes, formulations, parameters, percentages, quantities, characteristics, and other numerical values used in the specification and claims, are to be understood as being modified in all instancesRef. No.: JHU-42257.601 by the term “about” even though the term “about” may not expressly appear with the value, amount or range. Accordingly, unless indicated to the contrary, the numerical parameters set forth in the following specification and attached claims are not and need not be exact, but may be approximate and / or larger or smaller as desired, reflecting tolerances, conversion factors, rounding off, measurement error and the like, and other factors known to those of skill in the art depending on the desired properties sought to be obtained by the presently disclosed subject matter. For example, the term “about,” when referring to a value can be meant to encompass variations of, in some embodiments, ± 100% in some embodiments ± 50%, in some embodiments ± 20%, in some embodiments ± 10%, in some embodiments ± 5%, in some embodiments ±1%, in some embodiments ± 0.5%, and in some embodiments ± 0.1% from the specified amount, as such variations are appropriate to perform the disclosed methods or employ the disclosed compositions.

[0059] Further, the term “about” when used in connection with one or more numbers or numerical ranges, should be understood to refer to all such numbers, including all numbers in a range and modifies that range by extending the boundaries above and below the numerical values set forth. The recitation of numerical ranges by endpoints includes all numbers, e.g., whole integers, including fractions thereof, subsumed within that range (for example, the recitation of 1 to 5 includes 1, 2, 3, 4, and 5, as well as fractions thereof, e.g., 1.5, 2.25, 3.75, 4.1, and the like) and any range within that range.

[0060] As defined herein, “high-entropy alloys” (HEAs) are alloys containing 5 or more constituent elements, containing significant proportions, or substantially equimolar amounts, of each. As defined herein, “multi-principal element alloys” (MPEAs) are similar to HEAs but may include as few as two constituent elements, containing significant proportions, or substantially equimolar amounts, of each. Various HEAs and MPEAs exhibit one or more superior mechanical properties such as yield strength, fracture toughness, and fatigue resistance. Refractory MPEAs (RMPEAs) are MPEAs comprising refractory metals including Hf, Mo, Nb, Ta, Ti, V, W, and / or Zr. For the purposes of the present description, the metal alloys can be an MPEA comprising at least two elements selected from the group consisting of Hf, Mo, Nb, Ta, Ti, V, W, Rh, Zr, Li, Na, K, Rb, Cs, Mg, Ca, Sr, Ba, La, Ce, Al, Sn, Ga, In, Bi, Si, Ge, As, Sb, Cr, and Te. In some embodiments, the metal alloys can be a RMPEA comprising at least two elements selected from the group consisting of Hf, Mo,Ref. No.: JHU-42257.601Nb, Ta, Ti, V, W, and Zr. In some embodiments, the elements in a given alloy will possess varying binary mixing enthalpies with the other constituent elements (c.g., Ni-Co vs Ni-Fc or Co-Fe in CoFeNi alloys). Advantageously, the binary pairs with highly positive mixing enthalpies will also possess disparate oxygen interaction behavior.

[0061] The terms “base,” “bulk,” “base metal,” and “bulk metal” are used herein to describe the interior of the metal alloy component which is present prior to interstitial uptake and is substantially preserved during and after interstitial uptake as described herein.

[0062] The term “phase” is used herein to broadly describe a volumetric subdivision of a metal alloy component which possesses a resolvable crystal structure and range of compositions. Any given phase may share one, but generally not both, of these two attributes with any other phases present in the same metal alloy component. Prior to interstitial uptake, a metal alloy component may contain single or multiple phases. During and after interstitial uptake as described herein, the number of phases present increases due to phase separation of the original phases.

[0063] The term “precipitate” is used herein generally to encompass phase formation by nucleation and growth, as well as spinodal decomposition. This applies to: a single phase formed from one parent phase with or without consuming that entirety of the parent phase; a single phase formed from multiple parent phases with or without consuming that entirety of the parent phases; multiple phases formed from a single parent phase; or multiple phases formed from multiple parent phases.

[0064] The term “phase separation” is used herein to describe any instance of a homogenous phase transforming into two distinct phases. This can occur through a variety of phase transformation pathways both in the presence (i.e., spinodal decomposition) and in the absence (i.e., eutectoid transformations) of miscibility gaps. “Spinodal decomposition,” as well as “nucleation and growth,” are encompassed in this term.

[0065] The term “surface layer” is used herein to describe any portion of a microstructure which has undergone change, such as phase separation, due to interstitial uptake. This is not limited to the nanometer- scale layer comprising the physical surface, but rather extends to all transformed microstructural layers beneath this layer with the physical surface. This is also described herein by the term “near-surface region.”Ref. No.: JHU-42257.601

[0066] The term “interface” is used herein to describe the spatial region in which either (1) multiple phases intersect and adjoin, or (2) microstructural layers intersect and adjoin. These interfaces can be diffuse, or gradually changing in composition and / or structure across the interfacial region, such as the case with spinodal decomposition. These interfaces can also be sharp, or rapid in their compositional and / or structural changes across the interfacial region, as is the case with nucleation and growth.

[0067] The term “lamellae” and its related adjective “lamellar” are used herein to describe a resulting microstructure of thin, plate-like phases which alternate and repeat in a pattern.

[0068] The term “short-range order” is used herein to describe the preferential bonding coordination of atomic species to themselves or other species in a non-stochastic nature that is limited to the nearest several atomic neighbors, i.e., less than 1 nanometer in spatial extent for most metallic phases.

[0069] The term “medium-range order” is used herein to describe the preferential spatial concentration of atomic species with themselves or other species in a non-stochastic nature that is limited to more than the nearest several atomic neighbors but less than a distinctly resolved phase, i.e., typically between 1 and 5 nanometers for most metallic phases.

[0070] As used herein, “elevated temperature,” with respect to the methods described herein, corresponds to higher than room temperature but varying between alloy systems based upon diffusivity and interatomic enthalpic interactions, as readily understood by the person skilled in the art.

[0071] As used herein, “hardened” is a relative term, meaning that the surface subsequent to the treatment using the method described herein is harder after treatment than it was before treatment. Put another way, the metal alloy comprising interstitially-induced separated phases on or near a surface of a metal alloy is harder than the bulk metal alloy.

[0072] The term “porosity” is used herein to describe volumetric portions of the micro structure of the metal alloy component which, after applying the disclosed method, remain void of solid material. This can be caused by interstitially-induced phase separation which forms phases that melt, sublimate, or boil at the relevant temperatures of the disclosed method to leave behind such portions void of solid material.

[0073] Surface hardening, or case hardening, of metal components is widely used in manufacturing of articles or components made from steel, titanium, and other alloys. It isRef. No.: JHU-42257.601 most often accomplished through uniform surface interstitial uptake (i.e., B, C, N, O, H), surface precipitation of carbides, oxides or nitrides, selective surface phase evolution through surface heating and quenching, or coating of the surface. These methods achieve the end goal of creating a harder and more wear-resistant surface, but they do have shortcomings.

[0074] Uniform surface interstitial uptake methods produce hardened surfaces which have properties limited by the interstitial solubility of the bulk alloy composition. This leaves most alloy systems with only two variables to adjust to tune hardness: treatment temperature and the composition or concentration of the medium (solid, liquid, gas, or mixture) used to introduce the interstitials. The limited number of tuning variables inhibits both the number of alloys that can feasibly be case hardened and the extent to which an individual alloy can be case hardened. Moreover, there is a point at which the interstitial element concentration has reached the maximum solubility in the base alloy. At this point, precipitation of hard compound particles, such as carbides, oxides, or nitrides, begins. These precipitates are generally very hard but often serve as crack initiation sites and severely hinder the ductility and fatigue life of the article or component. This is caused by the hard particles’ interface with a now embrittled base alloy due to the maximum interstitial content. Another underlying cause of this can be lattice parameter mismatch.

[0075] Selective surface heating and cooling can be used to promote harder alloy phases at the surface than in the bulk of the article or component. A common example of this is case hardening of steel components by either quenching the surface at a higher rate or only heating the surface and proceeding with a quench. This results in very hard and brittle martensite which typically must then be partially tempered into austenite to improve toughness to acceptable levels. This approach limits the possible phases and requires tradeoffs in the resulting properties such as hardness, fatigue life, toughness, and corrosion resistance.

[0076] Coating is another common method for hardening the surface of metal components, as well as for thermal protection and environmental (corrosion and oxidation) protection. These coatings are typically atomized and sprayed onto the alloy surface, though they can be applied mechanically or by chemical / physical vapor depositions, as well. The coatings can be another metallic composition or ceramic. After application, the component is often heatedRef. No.: JHU-42257.601 to promote interdiffusion and bonding between the alloy substrate and the coating. The resulting adhesion of this step is highly dependent upon the lattice paramctcr / structurc of both the alloy substrate and the coating. This means that the bulk alloy composition and structure limits the number of compatible coatings, which may exclude the alloy from some applications.

[0077] Accordingly, there continues to exist a great need in the art to overcome the above problems with conventional alloy surface hardening and protection.

[0078] Many alloys possess miscibility gaps, which are regions in temperature-composition phase diagrams that result in separation into two or more distinct compositions, local atomic bonding structures, and / or crystal structures. A generic example of this phase diagram phenomenon is shown FIG. 1, though it is appreciated by the person skilled in the art that the phase transformation pathways can be much more complex. Miscibility gaps traditionally offer a pathway to control local chemistry throughout a metal component of a fixed bulk composition. Depending on the applied heat treatment, this can result in controlled bonding and diffusion environments including atomic short-range order, mediumrange order, long-range order, phase separation, and spinodal decomposition. Each of these conditions play an important role in immediate mechanical properties and the phase evolution that affects mechanical properties such as hardness. The types of miscibility gap- induced controlled bonding and diffusion all have important effects on solubility, uptake, and diffusion of interstitials. Provoking short to medium-range ordering can create localized regions with affinity, or aversion, to varying interstitial elements. This results in a “composite” behavior at fine length scales within single-phase regions experiencing locales of varying dislocation mobility and hardness. The resulting surface mechanical properties can be tailored solely through modulation of frequency and extent of atomic ordering, offering a “fine” tuning knob over traditional uniform interstitial uptake methods.

[0079] The short to medium-range order control of interstitial uptake also offers a mechanism by which the phase evolution of oxides, carbides, borides, nitrides, hydrides, and other interstitial-derived precipitates can be controlled thermodynamically, kinetically, and spatially. This can be achieved due to the alteration of ability to accommodate interstitial elements by introduction of short to medium-range order in the cationic species. For instance, a heat treatment can be applied to an article or component that utilizes proximity toRef. No.: JHU-42257.601 a miscibility gap to promote short to medium-range order and / or clustering. Exposure to interstitial elements can occur during this step and / or during a following treatment. Upon introduction of the interstitial species, they will encounter an altered energetic landscape for the most stable interstitial sites depending on the clustering and ordering of the cations. The heat treatment can then be extended, or another heat treatment with an inert gas or interstitial-element atmosphere can be performed, to promote preferential phase evolution. Alternatively, the article or component could undergo this treatment during normal service conditions. In summary, an alloy can be annealed at a temperature that introduces or suppresses ordering (or clustering) of desired metallic elements. Then interstitial element(s) are introduced, which are promoted or excluded in certain interstitial lattice sites due to the metallic ordering (or clustering). This creates a tunable intermediate step of selected local atomic chemistry and structure that could promote (or inhibit) desired (or undesired) interstitial-derived compounds. When compared to the example of martensitic case hardening of steel, this procedure offers additional pathways to promote preferred and suppress unwanted interstitial-derived phase precipitation with a greater control of precipitate location and concentration throughout the surface while offering an evolution pathway to previously unattainable phases. It also provides a hierarchy of phases available for formation depending upon the local concentration of the interstitials. This can allow for targeted selection and formation of interstitial-derived precipitates that offer enough hardness to the surface without posing a substantially high risk of crack initiation at the interface. The resulting atomic ordering and interstitial content in the alloy matrix around the precipitate can also be tuned to absorb the deformation energy that would otherwise cause crack initiation at the precipitate interface.

[0080] For the purposes of the present description, a “miscibility gap” covers not only nucleation and growth, e.g., as seen in the HINbTiZr, but also spinodal decomposition.

[0081] For the purposes of the present description, “heat treatment” or “heating” can be interpreted as heating the bulk sample or alloy to a uniform temperature by means of radiation, convection, induction, or resistive heating or localized surface heating, e.g., from a laser or electron beam.Ref. No.: JHU-42257.601

[0082] For the purposes of the present description, the interstitial element(s) comprise at least one of B, O, N, C, H, or any combination thereof. In some embodiments, at least one interstitial element is introduced as a gas, liquid, or a solid medium.

[0083] Some high-performance alloys including, but not limited to, refractory multiprincipal element alloys (RMPEAs), often exhibit miscibility gaps that cause phase separation into distinct compositions and / or crystal structures that possess vastly different interstitial solubilities. One mechanism of this phase separation is “spinodal decomposition” wherein specific chemical potential parameters lead to “uphill diffusion” with a negative diffusion coefficient. This typically results in a bicontinuous structure of two or more phases with different cation or interstitial content and, potentially, different bonding structures (see, e.g., FIG. 2 with off-base composition #1 and off-base composition #2). The composition of the resulting phases varies with time. Spinodal decomposition can be brought about by altering temperature which allows for solid solution instability during cooling, if the given chemistry and treatment temperature are above the spinodal line in the temperature dimension and aligned with the spinodal in the composition dimension. It can also be brought about isothermally and proceed prior to cooling due to compositional alterations that create solid solution instability at the given temperature when these compositional alterations shift the alloy into the spinodal region at the treatment temperature. The phase separation may also proceed by “nucleation and growth,” which entails a positive diffusion coefficient and “downhill diffusion” with a relatively constant precipitate concentration.

[0084] To this end, the presently disclosed subject matter utilizes the above concept but is unique in leveraging interstitial uptake to either stabilize low-temperature miscibility gaps or promote new miscibility gaps with respect to the interstitial elements at temperatures above the native interstitial-free alloy miscibility gap. This enables preferential tuning of phase evolution at the surface while preserving the chemistry and micro structure in the bulk.

[0085] The presently disclosed methods differ from conventional surface hardening both in microstructure and the potential for selective tuning of the phases that form in the surface layer. The phase separated microstructure with properly selected phases can produce a hardened surface layer or a hard but tough layer due to a “composite” effect of brittle and ductile phases with diffuse interfaces. These interfaces can offer similar effects to interstitial-derived precipitates but pose a very low risk of crack initiation. SpinodalRef. No.: JHU-42257.601 decomposition with varying interstitial content can be achieved through miscibility gaps between metal elements that result in different interstitial solubilities, or they can be achieved through miscibility gaps between metal and interstitial elements that cause “uphill” diffusion of the interstitial and / or metal elements. The presently disclosed interstitial- induced spinodal decomposition is unique and promising due to only the near-surface region undergoing this transition, in contrast to spinodal decomposition caused only by treatment temperature, as that method would produce a spatially homogeneous transformation throughout the bulk.

[0086] The presently disclosed subject matter applies to both alloys with interstitial-free miscibility gaps and alloys that produce miscibility gaps with respect to interstitial elements. It is comprised of a heat treatment in the presence of gas, liquid, or solid medium containing interstitial elements. The uptake of interstitials into the alloy upon heating can stabilize or introduce miscibility gaps solely at the near-surface region, which commences phase separation. Due to the nature of miscibility gaps, as well as spinodal decomposition, this phase formation behavior can be tuned by varying temperature, interstitial species, interstitial concentration, and pre-exposure metallic composition and ordering / clustering.

[0087] As will be described in Example 1 below, and is shown in FIG. 2, the presently disclosed methods were used to initiate near-surface spinodal decomposition during heat treatment of equiatomic NbTiZr in the presence of oxygen. The sample has a single-phase microstructure in the center while the surface underwent spinodal decomposition and a hierarchy of microstructure-refining phase separations caused by the influx of oxygen. These oxygen-driven phase separations covered the compositional range from dilute interstitial oxygen to suboxides and oxides. One phase in the initial spinodal decomposition layer was rich in Nb, slightly depleted in Zr and Ti, and low in O. The other phase in this layer had moderate levels of Zr and Ti, was depleted in Nb, and had higher O content. It also was observed that these two phases promoted the formation of hard suboxide and oxide compounds which further enhance hardness and wear-resistance. Miscibility gaps and spinodal decomposition can offer a tuning method for phase evolution. The multi-phase region with varying interstitial content and metal element composition will have multiple and possibly competing interstitial-derived compounds that can form with further heat treatment or service. This allows for selection and control of the specific chemistry andRef. No.: JHU-42257.601 structure of the compounds (such as oxides, carbides, et cetera) via control of the preceding spinodally-dccomposcd phases. The unique kinetics and energetics at these diffuse interfaces can also lead to formation of metastable yet beneficial compounds, such as complex oxides.

[0088] Accordingly, miscibility gaps can, either directly through miscibility-induced precipitation or indirectly through subsequent heating of interstitial rich or depleted zones, promote long-range ordered phases that include intermetallics. Typically, these types of phases are dispersed throughout the component, such as in jet turbine blades of nickel-based superalloys. The miscibility gap pathway offers selective precipitation at the surface while maintaining base alloy properties throughout the bulk of the component. Some service environments are so demanding on the available alloys that they require surface coatings to protect the component from oxidation, corrosion, or mechanical impact and wear. These coatings can be as simple as adhering to the component surface, or they can be diffusion coatings that cause coating elements to penetrate the surface and alloy elements to penetrate the coating when heat treated or in service. Any degree of miscibility-induced ordering or precipitation, from short-range atomic ordering to spinodal decomposition, can be tuned to promote phases of the proper lattice parameter and / or density to promote adequate coating adhesion. The interstitials selectively distributed by short to medium-range order in a singlephase surface can slow the diffusion of coating elements into the alloy and alloy elements into the coating, prolonging life of the coating and therefore the component. Interstitially- induced separated phase surfaces offer even higher control over the interdiffusion between the alloy and the coating, even enhancing interdiffusion of some elements while suppressing that of others. This will expand the number of compatible coatings for a given alloy, increasing the potential performance of the alloy in harsh environments. Alloy selection and heat treatment for proximity to miscibility gaps in temperature-composition space offers new and impactful pathways to improved surface hardening of metal alloy components, as detailed above.

[0089] Accordingly, in a first aspect, a method of preparing an interstitially-induced miscibility gap resulting in separation of phases on or near a surface of a metal alloy is described, said method comprising:Ref. No.: JHU-42257.601 heating a metal alloy to an elevated temperature that may introduce ordering or clustering of metallic elements of said metal alloy; and introducing at least one interstitial element at the elevated temperature, which is promoted or excluded in certain interstitial lattice sites due to the metallic ordering or clustering of the metallic elements of said metal alloy, wherein two or more different phases are separated due to interstitially-induced immiscibility, and wherein increased interstitial uptake at elevated temperature may promote a sequence of further phase separations of the initially-separated phases, resulting in a hierarchy of multiple phase separations as interstitial uptake proceeds.

[0090] In some embodiments, the separation of phases can proceed by either spinodal decomposition or nucleation and growth due to an interstitially-induced miscibility gap. In some embodiments, the metal alloy prior to heating is a single-phase or multi-phase metal alloy in which the heating causes phase decomposition of one or multiple phases present prior to heating. In some embodiments, the at least one interstitial element comprises O, N, C, B, H, or any combination thereof. In some embodiments, the at least one interstitial element is introduced as a gas, liquid, or a solid medium comprising the at least one interstitial element. In some embodiments, the metal alloy that is heated utilizes proximity to a native miscibility gap present in the interstitial-free alloy, or the formation of an entirely new miscibility gap due to the introduction of interstitial elements. In some embodiments, the interstitially-induced miscibility gap(s) and proceeding phase separation(s) result in a refined microstructure at or near the metal alloy surface, which is substantially maintained during subsequent phase transformations. In some embodiments, the metal alloy comprises at least two elements selected from the group consisting of Hf, Mo, Nb, Ta, Ti, V, W, Rh, Zr, Li, Na, K, Rb, Cs, Mg, Ca, Sr, Ba, La, Ce, Al, Sn, Ga, In, Bi, Si, Ge, As, Sb, Cr, and Te. In some embodiments, the metal alloy comprises at least two elements selected from the group consisting of Hf, Mo, Nb, Ta, Ti, V, W, and Zr.

[0091] In some embodiments, the interstitially-induced separated phases offer hardness to the surface of the metal alloy without posing a substantially high risk of crack initiation at an interface within the metal alloy. In some embodiments, a single-phase or multi-phase base, or interior, of the metal alloy is substantially structurally, chemically, and mechanically preserved during interstitially-induced phase preparation in the near-surface region. In otherRef. No.: JHU-42257.601 words, interstitially-induced phase separation is restricted to the near-surface region which substantially preserves the original state of the metal alloy base, or interior. In some embodiments, a single-phase or multi-phase base, or interior, of the metal alloy substantially maintains its ductility during interstitially-induced phase preparation in the near-surface region. In other words, interstitially-induced phase separation is restricted to the near- surface region which substantially preserves the ductility of the metal alloy base, or interior. In some embodiments, the interstitially-induced separated phases can be tuned by varying temperature, treatment time, interstitial species, interstitial concentration / flux, and preexposure metallic composition and ordering / clustering.

[0092] In some embodiments, the interstitially-induced separated phases are coated with a coating selected from an oxidation-resistant coating, a corrosion-resistant coating, a mechanical impact-resistant coating, a wear-resistant coating, or any combination thereof. In some embodiments, the method of preparing further comprises additional heat treatment, including localized surface heating, to promote preferential phase evolution. In some embodiments, the method of preparing further comprises additional heat treatment in the presence of an inert gas or additional interstitial elements to promote preferential phase evolution. In some embodiments, the method of preparing further comprises precipitating stable compounds onto or from the interstitially-induced separated phases, wherein the stable compounds comprise at least one of oxides, carbides, borides, hydrides, nitrides or any combination thereof, including sub-stoichiometric variations such as suboxides. In some embodiments, the method does not include anodization of the surface or microarc oxidation.

[0093] In some other embodiments, the base alloy further comprises at least one porosityinducing metal element, which during heating and interstitial uptake, forms a phase at the near-surface region (e.g., another metallic phase, an oxide, a carbide, a boride, a nitride, a hydride, or other interstitial-derived precipitate of said porosity-inducing metal element) that has low melting, boiling, or sublimation point and as such, at the temperatures of the method, the newly formed phase at the surface is volatile / melts / sublimates, inducing porosity at the surface. This can be particularly advantageous for biomedical implants. For example, the at least one porosity-inducing metal element can be Mo, which when the interstitial element is oxygen, MoOa is formed, which completely volatizes around I155°C.Ref. No.: JHU-42257.601Both liquid- and gas-induced porosity are possible through this mechanism. This allows for a wide range of controlled porosity structures in the near-surface region by simply adjusting the amount of the at least one porosity-inducing metal element (or a comparable low-melting or volatile phase former) content, oxidation temperature, and oxidation time. It should be appreciated by the person skilled in the art that the at least one porosity-inducing metal element can be any of the at least two elements of the metal alloy, as described herein, and will vary depending on the conditions of the method, on the interstitial element chosen and the other species present in the metal alloy, hence the tunability of the porosity. In some embodiments, the at least one porosity-inducing metal element is selected from Mo, W, and Cr.

[0094] In a second aspect, a metal alloy comprising interstitially-induced separated phases on or near a surface of a metal alloy is described. In some embodiments, the interstitially- induced spinodally separated phases are positioned between a single-phase or a multi-phase bulk alloy composition and at least one coating. In some embodiments, the metal alloy comprises at least two elements selected from the group consisting of Hf, Mo, Nb, Ta, Ti, V, W, Rh, Zr, Li, Na, K, Rb, Cs, Mg, Ca, Sr, Ba, La, Ce, Al, Sn, Ga, In, Bi, Si, Ge, As, Sb, Cr, and Te. In some embodiments, the metal alloy comprises at least two elements selected from the group consisting of Hf, Mo, Nb, Ta, Ti, V, W, and Zr. In some embodiments, the metal alloy comprises a MPEA. In some embodiments, the metal alloy comprises a RMPEA. In some embodiments, the metal alloy is neither a MPEA or a RMPEA. In some embodiments, the metal alloy comprises a phase- separated microstructure comprising a hard surface layer. In some embodiments, the metal alloy comprises a matrix-precipitate microstructure comprising a hard surface layer. In some embodiments, the metal alloy comprises a spinodally decomposed microstructure comprising a hard surface layer. In some embodiments, the metal alloy comprises a phase- separated microstructure comprising a surface layer comprising a composite of brittle and ductile phases with diffuse or sharp interfaces. In some embodiments, the metal alloy comprising interstitially-induced separated phases on or near a surface of a metal alloy is prepared according to the method of the first aspect. In some embodiments, the metal alloy comprises a surface layer with porosity formed by melted or sublimated / volatized metallic, oxide, carbide, hydride, boride, or nitride phases formed during the interstitially-induced phase separation. Accordingly, inRef. No.: JHU-42257.601 some embodiments, the hard surface layer comprises pores (i.e., has a porosity), induced as described herein.

[0095] In a third aspect, an article of manufacture comprises the metal alloy comprising interstitially-induced separated phases on or near a surface of a metal alloy, as described herein. Articles of manufacture include, but are not limited to, engines (e.g., jet engines), gas turbines for electricity generation and naval propulsion, hypersonic flight, fission nuclear reactor components, fusion nuclear reactor components, mining and drilling equipment, machining components (e.g., cutting tools), metal working components (e.g., rolls or forging dies), and biomedical implants (e.g., arthroplasty, dental implants, or stents).

[0096] Other aspects, features and embodiments of the invention will be more fully apparent from the ensuing disclosure and appended claims.EXAMPLES

[0097] The following Examples have been included to provide guidance to one of ordinary skill in the art for practicing representative embodiments of the presently disclosed subject matter. In light of the present disclosure and the general level of skill in the art, those of skill can appreciate that the following Examples are intended to be exemplary only and that numerous changes, modifications, and alterations can be employed without departing from the scope of the presently disclosed subject matter. The descriptions and specific examples that follow are only intended for the purposes of illustration, and are not to be construed as limiting in any manner.EXAMPLE 1Enhanced Hardness of Multiprincipal Element Alloys by Hierarchical Oxides1.1 Overview

[0098] Refractory multiprincipal element alloys (RMPEAs) offer superiority to incumbent high-temperature structural alloys due to high melting points and the retention of strength at elevated temperatures. Liu et al., 2022; Senkov et al., 2018a; Senkov et al., 2019; Senkov et al, 2018b; An et al., 2021; Zyka et al., 2019. Of this class of RMPEA alloys, those containing Group IV and V elements possess adequate ductility, low density, and the necessary formability. Liu et al., 2022; Senkov et al., 2018a; Senkov et al., 2019; Senkov etRef. No.: JHU-42257.601 al., 2018b; An et al., 2021 ; Zyka et al., 2019; Senkov et al., 2021 a; Ye et al., 2020; Lei et al., 2018; Lei ct al., 2020; Dong et al., 2021.

[0099] Efforts to improve the poor oxidation resistance with additions of passivating elements, however, leads to formation of brittle phases. Yang et al., 2017; Gorr et al., 2021; Butler et al., 2017; Butler et al., 2022; Lo et al., 2019; Lo et al., 2022. Chang et al., 2018. Group IV and V elements have vastly different oxygen solubilities and affinities, which convolutes the predictability of oxide formation and the ability to design an oxidation resistant alloy. In the present disclosure, it was discovered that the formation of hierarchical heterogeneities, from nanoscale structural ordering to coarsened tertiary phases, results in a tunable gradient subsurface hardness ranging from 9.8 GPa in the oxide-adjacent metal to a 601% increase with 23.1 GPa in the oxides. This mechanistic understanding of the oxygen influence on phase formation informs the design of new alloys in this compositional space to enhance oxidation resistance and obtain exceptional hardness while preserving the desired properties and mechanical integrity of the bulk alloy.

[0100] More particularly, the presently disclosed subject matter investigates equimolar NbTiZr to examine phase stability, oxide formation, and microstructure at elevated temperatures because of its low density of 6.63 g / cm3, high strength and good ductility up to 1200 °C, and the presence of its comprising elements in half of all RMPEA compositions studied. Senkov et al., 2018a; Senkov et al., 2019; Senkov et al, 2018b; Zyka et al., 2019.

[0101] Using high-resolution characterization, the presently disclosed subject matter uncovers how complex hierarchical microstructures evolve temporally and spatially during isothermal oxidation of NbTiZr at 1050 °C in atmospheric air. The novel microstructural evolution led to exceptional hardness with a maximum enhancement of 601% over the substantially interstitial-free alloy. Senkov et al., 2018b. Initial uptake of interstitial oxygen created structural inhomogeneities within the BCC solid solution base metal. Increased oxygen levels led to spinodal decomposition into Nb-rich and (Zr,O)-rich phases, which revealed the stabilization of a high-temperature miscibility gap at least 500 °C beyond the critical point of the O-free alloy. Kumar et al., 1994.

[0102] Continued oxidation led to a coarsening front at the interface between the base metal and spinodal phases, driven by the interfacial strain introduced through the BCCRef. No.: JHU-42257.601 to co displacive transformation. Critically, a continuous titanium oxide formed parallel to the surface along this coarsening front and offers an internal scale, which can be tuned to reduce oxygen diffusion into the bulk. The fine phases formed from spinodal decomposition developed into suboxides near the surface. In addition to the contribution to remarkable hardening, these subsurface hierarchical heterogeneities provide a template for potential pathways to minimize oxygen transport into the bulk alloy and promote surface scale adhesion. This could prove to be an attractive alternative to adding classical passivating elements, especially if it lowers the concentration of alloying additions required to form protective complex oxides or reduce defects in suboxides. Gorr et al., 2021; Lo et al., 2019; Lo et al., 2022.

[0103] Each of these microstructural transformations directly impact hardness (suboxides > spinodal > lamellae > base metal) and allow for a gradient tuning of surface- directed mechanical properties without the evident undermining of ductility in the singlephase base alloy. The broader applicability of this hardening phenomenon to other alloys and kinetic regimes was confirmed by producing a comparable hardness increase in equimolar HfNbTiZr (FIG. 3) through a similar hierarchical phase separation that proceeded by nucleation and growth, as well as other RMPEAs containing Hf, Nb, Ta, Ti, Zr, and Al. While the hierarchical microstructure reported here does not alone provide exceptional oxidation resistance, a path has been identified to limit oxygen transport from the surface into the bulk by tuning the multi-layered and hierarchical phase evolution to obtain desirable structure, chemistry, and morphology with minimal alteration of the base alloy.1.2 Materials and Methods1.2.1 Alloy Processing

[0104] The alloys were cast in an arc-melting furnace under partial argon atmosphere using high purity elements. A cigar-shaped mold in a water-cooled copper hearth was used. The ingots were flipped and re-melted 5 times to ensure the elements were distributed homogeneously. The ingots were then encapsulated in fused quartz under a partial Ar atmosphere and given a homogenization heat treatment at 1250 °C for 24 hours, followed by an ice water quench in which the quartz was broken immediately upon entering the water. After homogenization, an ingot section was encapsulated, treated at 1050 °C forRef. No.: JHU-42257.60124 hours, and then quenched to investigate structure and order by x-ray analysis. 5-mm cubes for the bulk oxidation experiment was sectioned from the homogenized ingots, polished to 1200 grit SiC sandpaper, and oxidized in atmospheric air at 1050 °C for 10 minutes, 30 minutes, 1 hour, and 3 hours and at 1250 °C for 30 minutes. The samples were removed from the furnace and “quenched” on a large aluminum heat sink held at room temperature. Oxidation of NbTiZr was repeated with an ice water quench following 3 hours of oxidation to qualitatively confirm that the micro structure and extent of coarsening was not dependent upon quenching rate. All cubes were then cross-sectioned and polished to 0.05-pm colloidal silica for SEM analysis and preparation of TEM liftout samples.1.2.2 Nanoindentation

[0105] Instrumented nanoindentation and the Oliver-Pharr method, Oliver and Pharr, 1992, were used to measure the elastic modulus and hardness values across the oxygen diffusion zone. The NanoBlitz 3D method was used on a KLA iNano system, outfitted with a diamond Berkovich tip, to generate more than 2800 indents in the specified area with a 5-pm spacing. A maximum load of 50 mN (corresponding to 415-820 nm indentation depths as shown in FIGs. 4A-4D) and a Poisson’s ratio of 0.3 were used for all measurements. X and Y dimensions (and indent counts) of the 5 pm spacing indentation areas for the NbTiZr “10 min” (FIG. 4A), 30 min (FIG. 4B), and 3-hour (FIG. 4C), and HfNbTiZr (FIG. 4D) micro structures were 400x100 pm2 (80x20), 300x60 pm2 (60x12), 800x100 pm2 (160x20), and 320x100 pm2 (64x20), respectively. The “10 min” microstructure area was gathered from an uncoarsened region in the 30 min sample due to identical micro structure of the small fraction of remaining spinodal decomposition areas in the 30 min sample. Hardness and modulus values were averaged across the direction parallel to the layer interfaces, such that one-dimensional hardness and modulus plots can be examined across interfacial boundaries. All nanoindentation for tests displayed in FIG. 5 were performed as above with varied dimensions, but at least 10 indents per averaged row. The NanoBlitz method was repeated for a 10-pm indent spacing in NbTiZr to ensure that the measured values did not change FIG. 6, confirming that there was no additional interaction between neighboring indentation sites for the 5-pm indent spacing. This allows for reliable use of the 5-pm indent spacing to obtain a better spatial resolution across the microstructuralRef. No.: JHU-42257.601 layers. For the NbTiZr 3-hour sample, datapoints with a modulus or hardness that deviated more than 30% from the average of the given x-valuc (x-axis is perpendicular to the surface) were omitted (and colored black in contour plots) to better represent the progression of mechanical properties across the layers without convolution from unparallel interfaces, uneven topography, surface contamination, or other confounding variables.1.2.3 Microscopy and Microanalysis

[0106] X-ray total scattering measurements were taken from approximately 1-mm thick sections of the ingot. Both samples were homogenized, and one of them was aged, as described above but for 18 hours rather than 24 hours. The measurements were performed at the NSLS-ii synchrotron facility at the 28-ID-l beamline using a Perkin Elmer detector. Since a pair distribution function (PDF) is a 10 projection of the interatomic atom-atom distance weighted by the scattering form-factor, it is a good probe for the average short-, medium- and long-range order. When focusing on low r-regions and comparing two related PDFs, one can extract changes in the local chemical environment. The measurements were done in a PDF mode for total scattering, where the 'camera length' of the detector was 21.2 cm, and the beamline radiation was 0.1665 A. Calibration was done using Ni powder. The total scattering patterns were then azimuthally integrated using 'pyFAI' package, Ashiotis et al., 2015, and reduced and Fourier transformed to a PDF using 'pdfgetx3'. Juhas et al., 2013.

[0107] The cross-sectioned and polished samples were imaged in a Thermo Scientific Helios G4 UC Scanning Electron Microscope (SEM). Regions of interest were targeted for Transmission Electron Microscopy (TEM) liftouts using a Ga Focused Ion Beam (FIB) in the same Helios microscope. TEM analysis, including Scanning Transmission Electron Microscopy (STEM) and Electron Energy Loss Spectroscopy (EELS), of Region II was performed in a Thermofisher Spectra 300kV fitted with an X- CFEG cold field emission source, panther STEM detection system, GIF Continuum 1066 EELS detector, Super-X symmetrical quad-crystal EDS system and 5th order corrected S- corr Cs probe corrector. STEM HAADF / BF imaging and Energy Dispersive X-ray Spectroscopy of Regions III and IV were conducted by the double- aberration-corrected JEOL ARM 300CF, equipped with dual 100-mm2SDD detectors operating at 300kV. The convergence semi-angle for STEM imaging and EDS was 25.6 mrad and the collectionRef. No.: JHU-42257.601 angles for STEM HAADF and BF imaging were 54 mrad, approximately 220 mrad, and approximately 14.2 mrad, respectively. S / TEM analysis of the spinodal structure was performed on a single-aberration-corrected JEOE GrandARM2 at an accelerating voltage of 300kV. Diffraction patterns were collected on a JEOE OneView detector with a camera length of 500 mm. STEM-EDS was collected with JEOL dual-EDS detectors. STEM-EDS of Region IV was performed on a JEOL F200 with JEOL dual-EDS detectors. High- resolution STEM images from Region III were filtered with the radial difference filter produced by HREM Research, Inc.1.3 Results and Discussion1.3.1 Order and Miscibility in NbTiZr

[0108] It was initially important to understand the initial deviations from perfect homogeneity due to thermal effects in the absence of oxygen to establish a baseline for comparison upon further investigation of oxygen-induced effects. Chemical short-range order (SRO) of the unoxidized single-phase BCC NbTiZr alloy was investigated to better understand the effect of thermal annealing on SRO in the absence of oxygen. Synchrotron x- ray total scattering experiments were performed followed by a pair distribution function (PDF) analysis of the “homogenized” (1250 °C, approximately 0.75 homologous temperature TH) and homogenized + aged (1050 °C, approximately 0.6 TH) conditions with no oxidation treatment, which represent the dilute limit of oxygen from processing impurities. The PDF results (FIG. 7A) show a tendency for SRO chemical clustering, as indicated by the increase in Zr-Zr first nearest-neighbor coordination at the expense of X-X (X=Nb,Ti).

[0109] To determine the influence of dissolved interstitial oxygen on SRO, electron diffraction was performed on the base metal adjacent to the spinodal layer. This revealed a single-phase BCC structure with evident diffuse scattering (FIG. 7C) that suggests fine real- space features experiencing a structural distortion due to chemical inhomogeneity without undergoing a structural phase transformation, as shown in BCC-co alloys. Sanchez et al., 1978; De Fontaine et al., 1971. This deviation from homogeneity is explained by the uptake of interstitial oxygen in this region, which influences the chemical coordination and structural symmetry as found in similar alloys with oxygen added during melting. InterstitialRef. No.: JHU-42257.601 complexes, such as these, have been found to significantly enhance mechanical properties such as hardness, tensile strength, ductility, and damping capacity. Ye ct al., 2020; Lei ct al., 2018; Lei et al., 2020; Ma and Wu, 2019. Nanoscale complexes are often precursors to phase formation, and in this case precede spinodal decomposition. The absence of diffuse scattering in diffraction of the spinodal phases, along with the evident peak splitting, confirm that the nanoclusters were the source of diffuse scattering in the base metal before spinodal decomposition when they separated into a disordered BCC phase with a distinct lattice parameter and diffraction peaks.

[0110] During oxidation, the single-phase base metal initially undergoes spinodal decomposition (FIG. 7B, FIG. 7D), followed by coarsening of the spinodal structure into aligned lamellae (FIG. 7E). The influence of oxygen on spinodal decomposition at the interior edge of the ODZ is demonstrated by the chemical segregation in STEM-EDS maps and electron diffraction peak splitting (FIG. 7C, FIG. 7D). Binodal phase separation and spinodal decomposition have not been observed in NbTiZr in previous studies due to kinetic limitations at the lower temperatures within the oxygen-free spinodal and binodal regions. Senkov et al., 2018b; Senkov et al., 2019; Whitfield et al., 2021; Kumar et al., 1994; Cao et al., 2022.

[0111] The observed spinodal decomposition during oxidation of NbTiZr at 1050 °C occurs for the following reasons: (1) interstitial oxygen alters thermodynamics to produce a high-temperature miscibility gap; (2) oxygen clusters into nanoscale complexes with Zr, inducing a higher amplitude of compositional fluctuation; and (3) the high-temperature isothermal oxidation process increases diffusivity of metal atoms relative to temperatures in the O-free miscibility gap, including that of sluggish Nb. The interstitial complexes and the following spinodal decomposition lay the kinetic and thermodynamic foundation for the hierarchy of phase evolutions with further oxygen uptake that dictate mechanical and oxidation properties.1.3.2 Oxygen-Induced Exceptional Hardness

[0112] The temporal series of isothermal oxidation (FIG. 7B-FIG. 7E, FIG. 7F) provided insight into both hardness and phase evolution, which provides a mechanism for control of a surface-directed gradient in hardness with a maximum value of 23.1 GPa. After 10 minutes (FIG. 7B), the ODZ had clear spinodal decomposition near the base metalRef. No.: JHU-42257.601 interface and a spinodal-like microstructure that lost bi-continuity near the surface. After 30 minutes (FIG. 7E), the spinodal structure at the base metal interface had mostly coarsened into a narrow band of lamellae with a few locations still remaining uncoarsened. This coarsening consumed the entire base metal interface and extended further towards the surface after 3 hours of oxidation (FIG. 7F). The room-temperature nanoindentation hardness prior to lamellar coarsening was >9.8 GPa in the single-phase base metal (6.7 at.%0), approximately 18 GPa in the spinodal microstructure, and averaged 22.6 GPa in the subsurface oxides (FIG. 8A). While the hardness range decreased to 7.6-12.4 GPa after 3 hours of oxidation due to the coarsening of the spinodal structure into lamellae, the elastic modulus of each region remained relatively unchanged and ranged between 103-153 GPa across the layers (FIG. 8C). Regions of different lamellae orientation produce local variations of the hardness and modulus within the layer. This drastic change in hardness and modulus between spinodal phases and lamellae is highlighted in nanoindentation of the coarsening front (FIG. 8B). For comparison, oxygen-free NbTiZr has an elastic modulus of 80.4 GPa and hardness of 3.3 GPa, Senkov et al., 2018b, while the C-103 alloy with approximately 9.5 at.%0 (16,700 weight ppm) has a reported hardness of about 1.9 GPa (190 HV), Sankar et al., 2013, and the hardness of monoclinic ZrO2 ranges from 9.8-13 GPa. Al-Khatatbeh et al., 2010. In addition to the expansion of the oxygen-induced hardness increase to equimolar HfNbTiZr (FIG. 3), the test matrix was expanded further to more compositions and treatment temperatures. These results are displayed in FIG. 5 and include: equimolar NbTiZr at 1250°C for 30 minutes (FIG. 5A); equimolar HfNbTaTiZr at 1250°C for 30 minutes (FIG. 5B); 32.33Nb-32.33Ti-32.33Zr-3Al (at.%) at 1250°C for 30 minutes (FIG. 5C); 39Nb-39Ti-19Zr-3Al at 1250°C for 30 minutes (FIG. 5D); 39Nb-39Ti-9.5Zr- 9.5Hf-3Al at 1050°C for 1 hour (FIG. 5E); and 39Nb-39Ti-9.5Zr-9.5Hf-3Al at 1250°C for 30 minutes (FIG. 5F). This wide range of compositions demonstrate the broader applicability to other alloys and the varied treatment times and temperatures show the tunability of results by varying processing parameters.

[0113] Five distinct microstructural transformations occurred during oxidation that give rise to a hierarchical microstructure. (Transformation 1): formation of oxygen and nitrogen interstitial complexes in the base alloy; (Transformation 2): rapid spinodal decomposition of the single- phase base alloy; (Transformation 3): separation of a thirdRef. No.: JHU-42257.601 phase upon increased oxygen content; (Transformation 4): conversion from metallic phases with dissolved oxygen to suboxidcs and monoxides; (Transformation 5): coarsening of the spinodal structure into lamellae, which includes the displacive BCC-to-co transformation and formation of an internal TiO scale. While further transformations will occur to complete the oxidation, our focus is the initial stages in these subsurface regions.

[0114] The base metal adjacent to the ODZ in the 30 min sample experiences a minimum average increase of 206% relative to the interstitial-free value, Senkov et al., 2018b, with the innermost x-distance averaging 10.1 GPa. This is directly attributed to the structural deviations from perfect symmetry and homogeneity induced by oxygen uptake in Transformation 1 (FIG. 7), which corresponds with interstitial bulk-dopant effects in similar alloys with lower levels of interstitial doping. Ye et al., 2020; Lei et al., 2018; Lei et al., 2020; Ma et al., 2019.

[0115] The spinodal decomposition (Transformation 2) provides a hardness increase of up to 463% over the interstitial-free alloy in the two-phase region with a maximum average of 18.5 GPa. Regions between the spinodal and surface are richer in oxygen and take on monoxide compositions and structures (Transformation 4) while maintaining the fine spinodal-like microstructure; this leads to a maximum hardness increase of 601% near the surface with the highest measurement of 23.1 and the highest x-distance average of 22.6 (585% increase). After further oxidation, the metallic spinodal phases coarsen into aligned lamellae, growing towards the surface (Transformation 5) and consuming the monoxide region. This coarsening reduces the hardness to values intermediate of the interstitial-rich base metal and the spinodal microstructure. FIGs. 8A-8C illustrate the gradient hardness and elastic modulus resulting from the hierarchical microstructure, which is tunable through simple processing parameter manipulation (e.g., temperature, time, partial pressure of oxygen and nitrogen).1.3.3 Effect of Oxygen on Phase Formation

[0116] The genesis of the interstitial complexes and spinodal decomposition were resolved by initial characterization, but the mechanism behind the coarsening and its propagation required further investigation. The Nb-rich, O-lean lamella exhibits the disordered BCC structure and the (Zr,O)-rich lamella is the hexagonal co phase, which isRef. No.: JHU-42257.601 well-known in Ti and Zr alloys (FIG. 9B). Sikka et al., 1982; Niinomi et al., 2016; Bailor et al., 2022; De Fontaine, 1988.

[0117] Geometric phase analysis (GPA) performed on a HAADF-STEM image of the BCC-co interface (FIG. 9C-FIG. 9D) reveals a misfit dislocation spacing ranging from 10-14 planes, corresponding to a semi-coherent misfit strain that varied between 7.1-10%. This diffusional co phase formation proceeds by collapse of BCC (111) planes and is favorable when P-stabilizers diffuse away from a location in Group IV P alloys, as occurs here with Nb due to the miscibility gap. Bailor et al., 2022; De Fontaine, 1988. This transformation entails larger cation diffusion distances than are involved in Transformations 1-4, causing it to occur last. The BCC-co transformation and the resulting reduction in coherency cause the spinodal structure to coarsen into lamellae to reduce the total interfacial energy.

[0118] Once the coarsening front initiates, it propagates toward the surface into the three-phase regions of suboxides and monoxides due to either semi-coherent strain relief or high mobility of incoherent boundaries. The discovery that oxygen flux and the co transformation promotes the coarsening reaction allows for control of the rate and degree of coarsening by selective doping to alter the interfacial strain and cation diffusivity, effectively tuning the desired gradient of surface-directed mechanical properties and the landscape for oxygen diffusion.

[0119] The hierarchical microstructure is fully formed after 3 hours of oxidation (FIGs. 10A-10C) with clearly defined regions resulting from the transformation events that succeeded the spinodal decomposition. All three oxidation exposures contained a region near the spinodal / lamellae interface with the base metal where the Ti-rich phase emerged, with the difference being whether enough time had passed for coarsening of the spinodal structure into lamellae (Transformation 5). After 3 hours, this third Ti phase presents as an approximately 10 nm sublamella that matches the hexagonal co-type structure (FIG. 10C).

[0120] STEM-HAADF and BF analysis of the nanotwinned Zr phase adjacent to Ti sublamella in Region III determined that this could be an orthorhombic suboxide of (Zr,Ti)CE Srilankite (Pbcn), Troitzsch et al., 2005, polar ZrCb (Pca21), Kersch and Falkowski, 2021, OI-ZrO2 (Pbca), Al-Khatatbeh and Lee, 2014, or a monoclinic suboxide of Baddeleyite Z1O2 (P21 / c), Al-Khatatbeh and Lee, 2014, (FIG. 10C). These nanoscaleRef. No.: JHU-42257.601 heterogeneities in a complex energetic landscape have important implications for both dislocation motion and the diffusivity of oxygen into the bulk. Ma and Wu, 2019; Wang ct al., 2022.

[0121] The phases between the coarsening front and the surface remain fine and globular in morphology (FIG. 10B). STEM-EDS reveals that the cations remain segregated and take on approximately 1:1 up to 1:2 stoichiometries with oxygen. Importantly, the fine size was retained despite obtaining higher oxidation states than the inner layers. This improves hardness and crack mitigation.

[0122] Oxidation of NbTiZr resulted in a hierarchical sequence of phase evolutions that manifested in distinct subsurface transformations and layers. This led to the first heterogeneity of nanometer- sc ale interstitial complexes which upon further oxidation progressed through spinodal decomposition, nanoscale suboxide phase separation, and a displacive transformation-induced coarsening reaction. These events resulted in an extensive oxygen diffusion zone that is not typically desirable in high-temperature service, though it did promote exceptional hardness with observed plasticity. Due to the enthalpic reactions and rapid kinetics of MPEAs with Group V elements, the extensive ODZ may be unavoidable without prior surface treatment. The ODZ layers did provide an altered substrate on which the surface oxides grow, providing a path to avoid common volume ratio issues between oxides and base alloys.

[0123] The reported RMPEAs with a mix of Group IV and V elements provide an opportunity to tune the phases formed in this hierarchical sequence by taking advantage of immiscibility, metastability, and interfacial strain during oxidation to both improve mechanical properties and reduce oxygen flux into the base alloy. This phenomenon extends to other MPEA composition spaces with moderate binary cation immiscibilities and comparable disparities in oxygen solubility and affinity. The kinetic barrier that has hindered spinodal decomposition in NbTiZr was overcome through the introduction of oxygen, which contributed to the layered phase formation hierarchy and refined micro structure. Further oxygen enrichment enhanced the stability of a displacive transformation by planar collapse of the Zr-rich BCC phase to the hexagonal co structure. This transformation was kinetically limited by the semi-coherent interfacial strain which acted to coarsen the phases into a lamellar morphology. While the omega phase exclusively formed as lamellae in theRef. No.: JHU-42257.601 equimolar NbTiZr, minor compositional changes can adjust morphology of the omega phase to prohibit lamellar formation such as in a Nb-Ti-Zr based alloy with reduced Zr and minor Al addition (see, e.g., FIGs. 11A-11C). In regions of equimolar NbTiZr with higher oxygen fluxes, the spinodal phases transitioned into defect-rich monoxides prior to the relatively sluggish coarsening. The mechanistic understanding of these transformations opens the path to tune crystal structure, defects, and morphology to alter the surface-directed mechanical properties by means of compositional modulation.EXAMPLE 2

[0124] Refractory alloys (containing high-melting elements such as Ti, V, Cr, Zr, Nb, Mo, Hf, Ta, W, etc.) have been suggested as potential replacements of incumbent implant alloys. Multi-principal elements alloys (MPEAs), which contain high proportions of all constituent elements and lack a single principal element, have seen the most focus within refractory alloy development for implants. Feng et. al, 2022. This is owed to their promising mechanical and corrosive properties, although the tradeoffs that apply to incumbent alloys have not yet been overcome. Alloys that are lean in Group IV elements (Ti, Zr, Hf) have high hardness and wear resistance but also have elevated elastic (Young’s) moduli. Alloys that are rich in Group IV elements behave similarly to Ti-based alloys and reduce stress mismatch at the expense of wear resistance. Senkov et. al, 2018a. This presents two pathways for optimization of implant viability for novel alloys: (1) reduce the modulus of Group IV-lean alloys without sacrificing wear resistance, or (2) increase the wear resistance of Group IV rich alloys without increasing the modulus or reducing biocompatibility.

[0125] Most recent developmental research has focused on the latter. This includes aqueous anodizing, microarc oxidation, and low temperature oxidation of well-established Ti alloys. The majority of developmental surface modification techniques, outside of rudimentary techniques such as sandblasting, have been focused on the principle of anodization. Anodized surfaces are formed during electrolysis of the surface in an acid, which forms a well-adhered TiCh surface oxide. Microarc oxidation takes this a step further by increasing the electrical potential to produce plasma arcs during anodization. This results in porosity in the surface oxides, enhancing osteoblast adhesion. Li et al., 2004. These methods allow for further modification by the common technique of plasma spraying toRef. No.: JHU-42257.601 deposit a hydroxyapatite coating, which improves osseointegration. A few studies have recently explored surface oxide formation from thermal oxidation rather than electrochemical anodization. They found that this process produces adequate TiCT formation on the surface of several Ti alloys when performed near 500°C, however, these surfaces lack significant advantage over anodized surfaces due to the general lack of nanoscale heterogeneities. Correa et al., 2022. This limitation can be overcome by alteration of alloy composition, as described herein, where a multilayered microstructure of oxides formed on NbTiZr and other alloys described herein after thermal oxidation, wherein each layer was composed of distinct morphologies of nanoscale suboxides.

[0126] Commercial and developmental alloys often possess miscibility gaps. This thermodynamic feature promotes separation of a single phase into two or more distinct phases when the proper local composition and temperature is met. This can lead to isostructural phases or phases of varying crystal structures, but the end result is always different compositions. These miscibility gaps typically manifest in the context of only metallic elements participating and the phase transformation being uniform throughout the bulk of the component. As disclosed herein, surface-directed gradient properties to optimize mechanical performance via miscibility gaps induced by interstitial element uptake can be created. In this example, adaptations of this method to create alloys and processing paths to synergistically optimize the desired properties of biomedical implants are described. This is demonstrated in Group IV-V refractory MPEAs using oxygen-induced miscibility gaps. It should however be appreciated by the skilled artisan that the underlying thermodynamics could apply to the right compositions of any class of alloys and using any interstitial element with proper interaction parameters that is absorbed from the surface during thermal exposure.

[0127] Equimolar NbTiZr prepared by plasma arc melting was cold-rolled to a reduction in height of 80%. A recrystallization anneal was not applied prior to the surface treatment with the intention to demonstrate the broader applicability of the treatment to a variety of starting microstructures. The heavily-deformed and un-recrystallized sample was then ground to 1200 grit SiC paper, followed by oxidation in ambient air at 1000°C for 10 minutes. The resulting structure mirrored the microstructural evolution described herein, with the main difference being an increase in titanium oxides around the simultaneouslyRef. No.: JHU-42257.601 recrystallizing grains near the coarsening front (FTG. 12). This adds additional microscale heterogeneities while still adhering to the overall hierarchical phase transformations caused by the oxygen-induced miscibility gap. In addition to the maximum hardness of 23.1 GPa near the surface of the 5mm cubic sample that was not cold-rolled and recrystallized (previously described), nanoindentation measurements from the center of this sample that confirm the preservation of the bulk mechanical properties that are favorable to structural implants have been obtained (FIG. 13, Table 1). The measured hardness was 3.46 ± 0.08 GPa and the modulus was 94.23 ± 0.91 GPa. This is nearly identical to the reported literature values of Vickers hardness of 3.33 ± 0.02 GPa (Senkov et. al, 2021b) and compressive Young’s modulus of 92 ± 3 GPa (Gong et. al, 2024) for interstitial-free NbTiZr. This confirms that the application of the method described herein can significantly increase surface hardness, and thereby wear resistance, while preserving bulk elastic properties closer to cortical bone than any incumbent implant alloy (Table 1). The result would be an implant with a longer lifetime due to reduced wear and stress shielding.Table 1: Hardness and elastic modulus values reported in literature and measured in this example (*). References are as follows: (a) Liu et al., 2004; (b) Odaira et al., 2022; (c) Chen et al., 2015; (d) Senkov et al., 2021b; (e) Gong et al., 2024; (1 Li et al., 2021.Ref. No.: JHU-42257.601

[0128] A second alloy, 33.3Nb-33.3Ti-8.3Mo-25Zr (at.%), was homogenized at 1250°C for 24hrs under vacuum and water quenched to eliminate casting microsegregations. It was then ground to 1200 grit and oxidized in ambient air at 1000°C for 10 minutes and 1050°C for 60 minutes. This alloy was selected because binary mixing enthalpies and oxygen affinities indicate that Mo will segregate to Nb-rich and Zr-lean phases during the hierarchical phase separation. As these phases undergo the transition to oxides upon further oxygen uptake, MoO will form. This oxide has low melting and boiling points. It melts at 795°C and completely sublimates around 1155°C. The vapor pressure at 900°C is IO-4Torr, in contrast to that pressure being obtained at 2200°C for ZrCh. Both liquid- and gas-induced porosity are possible through this mechanism. This allows for a wide range of controlled porosity structures in the oxide scale by simply adjusting Mo (or a comparable low-melting or volatile phase former) content, oxidation temperature, and oxidation time. The porosity formation was confirmed for both oxidation conditions by microstructural imaging in a scanning electron microscope, as shown in FIG. 14. This provides the opportunity to tailor the already advantageous microstructure even further to better promote osteoblast adhesion and osseointegration.

[0129] In some embodiments, the refractory multiprincipal element alloys described herein possess a reduced magnetic susceptibility and increased MRI compatibility relative to incumbent Ti and Co implant alloys.

[0130] In vitro studies with MC3T3-E1 preosteoblast cells from mice were performed to assess the biocompatibility of the alloys. Cylindrical discs were sectioned for the treated NbTiZr alloy as well as T1-6A1-4V, an alloy traditionally used in implants. The oxidized NbTiZr was then ground on diamond lapping paper to remove 50 pm of material from the surface and target a subsurface nanoscale oxide layer. The surface finishes were identical between the two groups, having been finished on 1 pm diamond lapping films. The samples were then incubated in a 96 well plate with preosteoblast cells for 24, 72, and 120 hours. Standard live / dead assay was used to identify live cells via the conversion of non- fluorescent calcein AM to highly fluorescent calcein through intracellular esterase activity, which is maintained within intact plasma membranes. DNA within dead cell nuclei with compromised cell membranes are stained red with ethidium homodimer- 1. Fluorescence intensity was quantitatively measured with a ClarioSTAR fluorescence microplate reader toRef. No.: JHU-42257.601 assess the relative viability and proliferation of cells in the presence of the sample (FIG. 15). Preliminary results showed significant improvement in cellular proliferation and viability in the NbTiZr compared to Ti-6A1-4V, with a 2x larger fold increase in live cell fluorescence over the course of the 5 day study. Dead cell fluorescence indicators slightly decreased over the 5 days for the NbTiZr while the Ti-6A1-4V saw up to a 2-fold increase in dead cell fluorescence by day 5. Fluorescence imaging (not shown) was then performed to qualitatively assess cellular adhesion and morphology on the surface of the alloy punches. Fluorescence imaging showed cells adhered to the surface of both alloy systems with traditional fibroblastic morphology consistent with the preosteoblast cell lines. While areas of the Ti-6A1-4V appeared to contain large areas of confluent cell populations, a lack of strong cell adherence was demonstrated by peeling edges of cell colonies that were freely floating away from the surface of the alloy. The NbTiZr samples did not exhibit any cells peeling off of the surface. Additionally, the surfaces of the NbTiZr samples had a more even distribution of cells over the entire surface that weren’t as densely packed when compared to the Ti-6A1-4V that were densely packed in some areas and barren in others. The in vitro studies act as a proof of concept and demonstrate the superior biocompatibility and osseointegration potential of the disclosed surface treatment on applicable alloys over the incumbent Ti-6A1-4V.Ref. No.: JHU-42257.601REFERENCES

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[0132] Although the foregoing subject matter has been described in some detail by way of illustration and example for purposes of clarity of understanding, it will be understood by those skilled in the art that certain changes and modifications can be practiced within the scope of the appended claims.

Claims

Ref. No.: JHU-42257.601THAT WHICH IS CLAIMED:

1. A method of preparing an interstitially-induced miscibility gap resulting in separation of phases on or near a surface of a metal alloy, said method comprising: heating a metal alloy to an elevated temperature that may introduce ordering or clustering of metallic elements of said metal alloy; and introducing at least one interstitial element at the elevated temperature, which is promoted or excluded in certain interstitial lattice sites due to the metallic ordering or clustering of the metallic elements of said metal alloy, wherein two or more different phases are separated due to interstitially-induced immiscibility, and wherein increased interstitial uptake at elevated temperature may promote a sequence of further phase separations of the initially- separated phases, resulting in a hierarchy of multiple phase separations as interstitial uptake proceeds.

2. The method of claim 1, wherein the separation of phases can proceed by either spinodal decomposition or nucleation and growth due to an interstitially-induced miscibility gap.

3. The method of claim 1, wherein the metal alloy prior to heating is a single-phase or multiphase metal alloy in which the heating causes phase decomposition of one or multiple phases present prior to heating.

4. The method of claim 1, wherein the at least one interstitial element comprises O, N, C, B, H, or any combination thereof.

5. The method of any of the preceding claims, wherein at least one interstitial element is introduced as a gas, liquid, or a solid medium comprising the at least one interstitial element.

6. The method of any of the preceding claims, wherein the metal alloy that is heated utilizes proximity to a native interstitial-free miscibility gap or the formation of an entirely new miscibility gap due to the introduction of interstitial elements.Ref. No.: JHU-42257.6017. The method of any of the preceding claims, wherein the interstitially-induced miscibility gap(s) and proceeding phase scparation(s) result in a refined microstructure at or near the metal alloy surface, which is substantially maintained during subsequent phase transformations.

8. The method of any of claims 1-7, wherein the metal alloy comprises at least two elements selected from the group consisting of Hf, Mo, Nb, Ta, Ti, V, W, Rh, Zr, Li, Na, K, Rb, Cs, Mg, Ca, Sr, Ba, La, Ce, Al, Sn, Ga, In, Bi, Si, Ge, As, Sb, Cr, and Te.

9. The method of any of claims 1-7, wherein the metal alloy comprises at least two elements selected from the group consisting of Hf, Mo, Nb, Ta, Ti, V, W, and Zr.

10. The method of any of the preceding claims, wherein the interstitially-induced separated phases offer hardness to the surface of the metal alloy without posing a substantially high risk of crack initiation at an interface within the metal alloy.

11. The method of any of the preceding claims, wherein a single-phase or multi-phase base, or interior, of the metal alloy is substantially structurally, chemically, and mechanically preserved during interstitially-induced phase preparation in the near-surface region.

12. The method of any of the preceding claims, wherein a single-phase or multi-phase base, or interior, of the metal alloy substantially maintains its ductility during interstitially-induced phase preparation in the near-surface region.

13. The method of any of the preceding claims, wherein the interstitially-induced separated phases can be tuned by varying temperature, treatment time, interstitial species, interstitial concentration / flux, and pre-exposure metallic composition and ordering / clustering.

14. The method of any of the preceding claims, further comprising coating the interstitially- induced separated phases with a coating selected from an oxidation-resistant coating, a corrosion-resistant coating, a mechanical impact-resistant coating, a wear-resistant coating, or any combination thereof.Ref. No.: JHU-42257.60115. The method of any of the preceding claims, further comprising additional heat treatment to promote preferential phase evolution.

16. The method of any of the preceding claims, further comprising additional heat treatment, including localized surface heating, in the presence of an inert gas or additional interstitial elements to promote preferential phase evolution.

17. The method of any of the preceding claims, further comprising precipitating stable compounds onto or from the interstitially-induced separated phases, wherein the stable compounds comprise at least one of oxides, carbides, borides, hydrides, nitrides or any combination thereof.

18. The method of claim 17, wherein the stable compounds include sub- stoichiometric variations such as suboxides.

19. The method of any of the preceding claims, wherein the metal alloy comprises a near- surface layer with porosity formed by melted or sublimated / volatized metallic, oxide, carbide, hydride, boride, or nitride phases formed during the interstially-induced phase separation.

20. A metal alloy comprising interstitially-induced separated phases prepared according to the method of any of claims 1-19.

21. The metal alloy of claim 20, wherein the interstitially-induced separated phases are positioned between a single-phase or a multi-phase bulk alloy composition and at least one coating.

22. The metal alloy of claims 20 or 21, wherein the metal alloy comprises at least two elements selected from the group consisting of Hf, Mo, Nb, Ta, Ti, V, W, Rh, Zr, Li, Na, K, Rb, Cs, Mg, Ca, Sr, Ba, La, Ce, Al, Sn, Ga, In, Bi, Si, Ge, As, Sb, Cr, and Te.Ref. No.: JHU-42257.60123. The metal alloy of any of claims 20-22, wherein the metal alloy comprises at least two elements selected from the group consisting of Hf, Mo, Nb, Ta, Ti, V, W, and Zr.

24. The metal alloy of any of claims 20-23, wherein the metal alloy comprises a phase-separated microstructure comprising a hard surface layer.

25. The metal alloy of any of claims 20-23, wherein the metal alloy comprises a phase- separated microstructure comprising a surface layer comprising a composite of brittle and ductile phases with diffuse or sharp interfaces.

26. The metal alloy of any of claims 20-23, wherein the metal alloy comprises a spinodally decomposed microstructure comprising a hard surface layer.

27. The metal alloy of any of claims 20-23, wherein the metal alloy comprises a matrixprecipitate microstructure comprising a hard surface layer.

28. The metal alloy of any of claims 20-23, wherein the metal alloy comprises a surface layer comprising pores.

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