Iron-tungsten redox system and applications of same

Abstract

This invention discloses a redox system comprising a Fe-xW powder architecture fabricated to possess excellent resistance to degradation at high temperatures during redox cycling between oxidized and reduced states, wherein x represents a concentration (at%) of tungsten (W) in the powder architecture and is in a range of about 1-75 at%, wherein the powder architecture comprises powder bed, powder suspension, foams, fibers, and / or printed microlattices.

Description

[0001]Attorney Docket No.: 616146.100555 IRON-TUNGSTEN REDOX SYSTEM AND APPLICATIONS OF SAME STATEMENT AS TO RIGHTS UNDER FEDERALLY-SPONSORED RESEARCH This invention was made with government support under CMMI-2015641 awarded by the National Science Foundation. The government has certain rights in the invention. CROSS-REFERENCE TO RELATED PATENT APPLICATIONS This application claims priority to and the benefit of U.S. Provisional Application Serial No. 63 / 619,401, filed January 10, 2024, which is incorporated herein in its entirety by reference. This application is also a continuation-in-part application of U.S. Patent Application Serial No. 18 / 367,019, filed September 12, 2023, which itself claims priority to and the benefit of U.S. Provisional Application Serial No.63 / 406,320, filed September 14, 2022, which are incorporated herein in their entireties by reference. FIELD OF THE INVENTION The present invention generally relates to the material science, particularly to an iron-tungsten redox system and applications of the same. BACKGROUND OF THE INVENTION The background description provided herein is to present the context of the invention generally. The subject matter discussed in the background of the invention section should not be assumed to be prior art merely due to its mention in the background of the invention section. Similarly, a problem mentioned in the background of the invention section or associated with the subject matter of the background of the invention section should not be assumed to have been previously recognized in the prior art. The subject matter in the background of the invention section merely represents different approaches, which in and of themselves may also be inventions. Work of the presently named inventors, to the extent it is described in the background of the invention section, as well as aspects of the description that may not otherwise qualify as prior art at the time of filing, are neither expressly nor impliedly admitted as prior art against the invention. Iron is an attractive energy-storage material due to its abundance, low cost, and non-toxicity. At high temperature, iron can be oxidized by steam and its oxides can be reduced by hydrogen, making iron a viable energy-storage material for the nascent rechargeable oxide battery (ROB) space. The ROB devices utilize the energy-storage material to convert a fixed amount of steam into hydrogen Attorney Docket No.: 616146.100555 which is then fed into a connected reversible solid oxide cell. Similar gas-conversion technologies could use iron as a low-cost, high-utility material, such as (i) chemical looping combustion, wherein iron oxide is reduced by a hydrocarbon fuel such as CH4, followed by a re-oxidation in air to produce heat, or re-oxidation with steam to form H2, (ii) carbon utilization, wherein CO2oxidizes Fe to produce CO as a chemical feedstock (e.g., for acetic acid), resulting in carbon sequestration and value-added CO2 use, or (iii) solar thermal redox, where extreme temperatures rather than reducing gases cause thermal reduction. All these technologies rely on the ability of the cycling material (e.g., iron) to convert from metal to oxide and vice versa while maintaining the same reactivity over multiple redox cycles, and to complete each half cycle in a practical timescale. While beds of pure Fe powders show promise during the first few redox cycles at the high temperatures (above 700 °C) necessary to achieve fast kinetics, they quickly lose reactivity due to powder sintering in both the metallic (Fe, Tm = 1538 °C) and oxide (Fe3O4,Tm = 1597 °C) states. Sintering greatly slows redox kinetics by (i) reducing the active surface area of the material via densification of the powders within the bed and (ii) blocking gas access to the metal powders via formation of a dense layer on the outer surface of the powder bed. Several strategies have been used to address this sintering-induced degradation. Frequently, the iron material is supported on an inactive oxide material such as titania, alumina, zirconia, or Mg-Al-O spinel. The underlying oxide support acts as a barrier to diffusion and sintering by preventing contact between nearby iron particles. The impact of porosity in these systems is key, as the materials with the best long-term performance, such as the FeTiO3materials, show stable or even increasing porosity with continued cycling, rather than densification of pores. Alloying iron with other metals is another strategy, with the candidate additions falling into three categories: inactive metals, inactive oxides, or redox active metals. Inactive metals - such as Ni and Cu - are not oxidized by steam under standard redox conditions, and instead provide a mechanically stable, metallic backbone during cycling. The interface between the metallic backbone and the iron oxide layer forming during oxidation helps accelerate reduction kinetics. For Fe-soluble inactive metals like Ni, the backbone provides a sink for Fe during reduction which helps stabilize the microstructure. Metals such as Al or Cr oxidize under steam into stable oxides which are not reduced by H2, effectively turning the alloy into the inactive oxide case discussed above. Redox-active metals comprise two main members, Mo and W. These refractory metals have very high melting point and thus provide sintering inhibition similar to that seen for oxide supports, while also contributing to the redox activity of the material. Our previous study of Fe-Mo redox cycling shows that Mo strongly inhibits degradation by sintering as compared to Fe-Ni. Some Attorney Docket No.: 616146.100555 degradation remains however: (i) Mo-rich phases segregate away from the surface of the microstructure as redox cycles are accumulated, and Mo loses its beneficial impact on the remaining Fe, and (ii) volatilization of MoO2(OH)2 causes further loss of Mo from the material due to chemical vapor transport. While initially observed to be detrimental to steam-hydrogen cycling characteristics, W has recently been shown to have a positive impact on the stability of iron undergoing high temperature redox cycling in different gas environments. However, a detailed examination of the microstructural evolution of high W-content Fe-W materials has not yet been conducted, nor has there been a study of this material under H2 / H2O cycling. Further, the use of micron- rather than nano-size powders in foams with mass of about 1 g is of practical interest and can strongly affect the structure and agglomeration of the powders. Any of these composition-based strategies can be combined with a structural approach besides a simple powder bed; for example, 3-D architectures (such as ink-printed lattices, extruded grates, and freeze cast foams) can provide a pore structure that lowers tortuosity, so that gas can easily enter and escape the reacting material. Our previous work with directionally freeze-cast foams shows that the directional channels and low tortuosity are effective for solid-gas reactions. Additionally, freeze casting is a low-cost manufacturing method that provides very fine pore structure, with channel widths as narrow as 20 µm. This is sufficiently wide to provide low tortuosity for hydrogen and steam, without allocating excessive foam volume to pores. However, freeze-cast foams undergo a unique degradation mechanism: due to the high aspect ratio of the lamellae, buckling is easily induced by the high stresses due to cyclic expansion and contraction during oxidation and reduction, respectively. For foams without a sintering inhibitor, sintering and densification occurs due to contact between neighboring lamellae that buckle and contact each other, eliminating the benefits of the foam architecture. Therefore, a heretofore unaddressed need exists in the art to address the aforementioned deficiencies and inadequacies. SUMMARY OF THE INVENTION In view of the above noted deficiencies and inadequacies, this invention provides a redox system comprising a Fe-xW (at%) (x being a concentration of 1-75 at% W) powder architecture (porous body) including but not limited to static powder bed, foam, fluidized powder bed, fibers, or powder suspension, which is fabricated to possess excellent resistance to degradation at a high- temperature during alternating oxidizing and reducing gas conditions. Oxidizing gases are those that Attorney Docket No.: 616146.100555 increase the oxidation state of Fe and / or W, including but not limited to H2O, O2, and CO2. Reducing gases are those that decrease the oxidation state of Fe and / or W, including but not limited to H2, CH4 (or other hydrocarbons including ethane, propane, butane, etc.), NH3, and CO. Oxidizing and reducing gas conditions may include mixtures of these gases, including additional inert component gases, including but not limited to Ar and N2. Gases including but not limited to: reducing (e.g., H2, CH4, N2), oxidizing: (e.g., O2, H2O, CO2), combined with any inert gases (e.g., N2, Ar). Redox cycling conditions include a temperature dependence, with excellent resistance to degradation between 500 – 1500 °C. The composition of the porous body shows excellent resistance to degradation in the range of 1- 75 at% W, preferably 5-33 at% W, with optimal results shown at 25 at% W and performance declining, but still showing better resistance to degradation than unmodified Fe. Additional components used in conjunction with W also show excellent resistance to degradation, including but not limited to Ni, Co, Cu and Mo, in the compositional ranges Fe-(5-33)(Mo+W), Fe-(5-33)(Ni+W), or Fe-(5- 33)(Mo+Ni+W), all in at%. In one aspect of the invention, the redox system comprises a Fe-xW powder architecture fabricated to possess excellent resistance to degradation at a high-temperature during redox cycling between oxidized and reduced states, wherein x represents a concentration (at%) of tungsten (W) in the powder architecture and is in a range of about 1-75 at%, wherein the powder architecture comprises powder bed, powder suspension, foams, fibers, and / or printed microlattice. In one embodiment, the redox cycling is performed with an oxidizing gas and an reducing gas alternatively, wherein the oxidizing gas are those that increase the oxidation state of iron (Fe) and / or tungsten (W) and comprises H2O, CO2, and / or O2, and wherein the reducing gas are those that decrease the oxidation state of Fe and / or W and comprises H2, hydrocarbons including CH4, ethane, propane and / or butane, NH3, and / or CO. In one embodiment, the oxidizing gas and the reducing gas are combined with any inert gases. In one embodiment, the high-temperature is in a range of about 500 – 1500 °C. In one embodiment, the powder architecture possesses the excellent resistance to degradation at about 800 °C during the redox cycling between the metallic and oxide states, with good reaction kinetics maintained up to a minimum of 100 full cycles. In one embodiment, the powder architecture is of Fe-1W, Fe-5W, Fe-10W, Fe-18W, Fe-20W, Fe-25W, Fe-30W, Fe-33W, Fe-40W, Fe-45W, Fe-50W, or Fe-75W. In one embodiment, the Fe-10W powder architecture with a lower volume fraction of W- containing phases of λ-Fe2W and FeWO4 shows lamellar densification and core-shell structure formation, due to Fe outward diffusion during oxidation. Attorney Docket No.: 616146.100555 In one embodiment, each of the Fe-18W and Fe-25W powder architectures comprises a sufficient volume fraction of W-rich phases including λ-Fe2W to inhibit sintering for α-Fe in the reduced state and FeWO4 to inhibit sintering for Fe3O4 in the oxidized state, thereby forming microstructures comprising two phases of Fe / λ-Fe2W and Fe3O4 / FeWO4, respectively. In one embodiment, each ligament in the Fe-25W powder architecture initially includes a mixture of bcc α-Fe(W) and µ-Fe7W6; during oxidation, α-Fe oxidizes to Fe3O4 and µ-Fe7W6 oxidizes to FeWO4and minor amounts of Fe3O4; and during reduction, the Fe3O4reduces back to bcc Fe, and the FeWO4 reduces more slowly to Fe2W. In one embodiment, there a mismatch in the stoichiometric ratio between Fe and W in moving from FeWO4(Fe:W=1) to Fe2W (Fe:W=2), whereby Fe from the surrounding matrix from prior reduction of Fe3O4is used to form the intermetallic Fe2W. In one embodiment, the second oxidation is, on average, about 40-80% faster than the first oxidation, and the second reduction is, on average, about 10-50% faster than the first reduction. In one embodiment, in the reduced state, the powder architecture includes powders which exhibit 2-10 µm ligaments interpenetrated by open porosity of the same size, wherein each ligament, in turn, includes an α-Fe(-1.6W) backbone decorated, and at times interrupted, with submicron (1-900 nm) λ-Fe2W grains and submicron CVT pores, and wherein this same reduced microstructure is maintained through 100 cycles with the micron-scale open porosity increasing with cycle number. In one embodiment, the oxidized state includes two interpenetrating networks of Fe3O4 and FeWO4. In one embodiment, the powder architecture, after the first cycle, comprises three levels of pore sizes including macroscopic freeze-cast channels (level I) separating lamellae which are distributed into colonies with aligned radial orientation; microscale sintering-inhibition pore network (level II) within each lamella, fully continuous, formed due to the sintering inhibition of W; and submicron chemical vapor transport (CVT) pores (level III), formed within each W-containing area in the lamellae, regenerated at each reduction half-cycle, thereby creating a hierarchical porosity structure in the powder architecture. In one embodiment, the microscale sintering-inhibition pores increase the surface area and thus reactivity of the ligaments and the submicron CVT pores greatly increase surface area and limit damage accumulation by effectively resetting the microstructure after each cycle, producing nanocrystalline λ– Fe2W with corresponding nanopores. In one embodiment, the microstructural stability allows for the powder architecture to grow progressively more porous (rather than denser) with cycling, and the W-rich phases become both Attorney Docket No.: 616146.100555 smaller in morphology and more homogeneously distributed throughout the ligaments, thereby increasing the sintering inhibition effect and generating more submicron pores to accelerate subsequent reaction. In one embodiment, the resistance is attributed to the sintering inhibition of W combined with the CVT mechanism of reduction. In one embodiment, the resistance is attributed to both the sintering inhibition inherent to W, and the nature of the intermetallic and mixed oxide compounds formed between Fe and W, such that W does not segregate and remains atomically mixed with Fe at all times. In one embodiment, W is atomically mixed with Fe, either in the intermetallic λ–Fe2W or in the mixed oxide FeWO4phases, thereby effectively preventing W from segregating into a pure W phase that reduces the degradation resistance of the Fe-25W powder architecture. In one embodiment, increasing W concentration has a minor effect on the oxidation rate but significantly extends the initial reduction time in the first cycle due to the slower reduction of more stable FeWO4, as compared to Fe3O4. In one embodiment, the chemical vapor transport (CVT) mechanism during reduction refines stable W-rich phases to a nanocrystalline scale, boosting the reaction kinetics of subsequent redox cycles. The reduction time could be adjusted by designing W concentrations and associated phase fractions of FeWO4. In one embodiment, higher W concentration enhances the degradation resistance of Fe-W lamellar foams during high-temperature redox cycling, resulting in a more stable lamellae structure. In one embodiment, the powder architecture further comprises Ni, Co, Cu, and / or Mo, in compositional ranges Fe-(5-33)(Mo+W), Fe-(5-33)(Ni+W), or Fe-(5-33)(Mo+Ni+W), all in at%. In one embodiment, the powder architecture is of Fe-(5-33)(Mo, W), wherein Mo and W are both active in oxidation and reduction. In one embodiment, the powder architecture is of Fe-9Mo-9W. In one embodiment, the powder architecture is of Fe-(5-33)(Ni, W), wherein W is active in oxidation and reduction while Ni is inert. In one embodiment, the powder architecture is of Fe-20Ni-20W. In one embodiment, the Fe-20Ni-20W powder architecture, after initial reduction, include a two-phase mixture of µ-Fe7W6and γ-Fe(Ni,W), with significant microporosity due to the sintering inhibition of W, and wherein during oxidation at 800 °C, these phases are oxidized to a three-phase mixture of (Fe,Ni)WO4, Fe3O4, and γ-Ni(Fe); upon subsequent reduction by H2, the powder architecture returns to their initial composition. Attorney Docket No.: 616146.100555 In another aspect, the invention relates to a device comprising the redox system as disclosed above. In a further aspect, the invention relates to a method for fabricating a redox system comprising milling a blend of an iron (Fe) or Fe-containing powder and a tungsten (W) or W-containing powder into a homogenous powder; and heat treating the powder blend to yield a powder architecture of Fe- xW, wherein x represents a concentration (at%) of W in the powder architecture and is in a range of about 1-75 at%. In one embodiment, the Fe or Fe-containing powder comprises a Fe powder, a Fe-rich powder, a Fe-containing oxide powder, a Fe-containing nitride powder, and / or a Fe-containing carbide powder. The W or W-containing powder comprises a W powder, a W-rich powder, a W-containing oxide powder, a W-containing nitride powder, a W-containing carbide powder, and / or ammonium paratungstate. In one embodiment, the powder architecture is of Fe-1W, Fe-5W, Fe-10W, Fe-18W, Fe-20W, Fe-25W, Fe-30W, Fe-33W, Fe-40W, Fe-45W, Fe-50W, or Fe-75W. In one embodiment, the powder architecture possesses the excellent resistance to degradation at about 800 °C during redox cycling between oxidized and reduced states, with good reaction kinetics maintained up to a minimum of 100 full cycles. In one aspect, the invention relates to a method for fabricating a redox system comprising providing a Fe-xW alloy; atomizing the alloy to form a powder; and optionally heat treating the powder to yield a powder architecture of Fe-xW, wherein x represents a concentration (at%) of W in the powder architecture and is in a range of about 1-75 at%. In one embodiment, said providing the Fe-xW alloy is performed by melting iron and ferrotungsten. In another aspect, the invention relates to a method for fabricating a redox system comprising producing a Fe-and W-rich oxide by sintering of precursor powders; pulverizing the Fe-and W-rich oxide to form a powder with Fe-xW oxide composition; and optionally heat treating the powder to yield a powder architecture with metallic Fe-xW composition, wherein x represents a concentration (at%) of W in the powder architecture and is in a range of about 1-75 at%. These and other aspects of the present invention will become apparent from the following description of the preferred embodiment taken in conjunction with the following drawings, although variations and modifications therein may be affected without departing from the spirit and scope of the novel concepts of the invention. Attorney Docket No.: 616146.100555 BRIEF DESCRIPTION OF THE DRAWINGS The accompanying drawings illustrate one or more embodiments of the invention and together with the written description, serve to explain the principles of the invention. Wherever possible, the same reference numbers are used throughout the drawings to refer to the same or like elements of an embodiment. FIG.1. (a) isothermal section of the Fe-W-O ternary phase diagram at 800 °C. The blue arrow indicates the overall phase evolution path for the Fe-25W alloy atomic ratio Fe:W=3, and the green phase boundary (Fe3O4 – FeWO4) indicates the fully oxidized composition. (b) Baur-Glässner diagram showing expected phase formation as a function of temperature and H2 partial pressure. In addition to Fe, FeO, and Fe3O4, thermodynamic data for the Fe-25W system are added in purple for comparison. (c) Arrhenius plot showing the various redox couples as a function of temperature and oxygen partial pressure (pO2). Feasible systems should fall within the 1:4 and 4:1 H2:H2O boundaries (red dotted lines); the Fe-25W system meets this criteria. FIG.2. SEM micrographs of the surface of a representative lamella in a Fe-25W foam (a) after initial reduction and sintering at 1200 ºC and (b) after 1 redox cycle at 800 ºC. Blue arrows indicate µ- Fe7W6, red arrows indicate α–Fe, and yellow arrows indicate λ–Fe2W, as determined by Z-contrast, with overall composition confirmed by XRD. (c) XRD spectra of Fe-25W freeze cast foams after initial reduction and sintering at 1200 ºC (black) and after the first redox cycle at 800 ºC (red). Before cycling, the foam includes a mixture of crystalline α-Fe (blue circle) and µ-Fe7W6 (green square), with some residual Fe3W3C carbide (purple diamond) from binder burnout. After cycling, the foam is returned to the equilibrium phases at 800 °C: α-Fe and λ-Fe2W (orange triangle), with the carbide removed by carbon oxidation. The broad hump between 40 and 48° is indicative of nano-crystalline λ- Fe2W formed by CVT reduction. FIG.3. (a) schematic depiction of the three hierarchical levels of porosity present in the foams, with inset (black circle) showing schematically the submicron pores. (b)-(d) Micrographs illustrating the hierarchical porosity present in reduced Fe-25W foams after the first cycle including three levels: (b) level I - macroscopic freeze cast channels, shown in a radial cross-section micrograph, (c) level II - microscale sintering-inhibition pore network, shown in surface electron image, and (d) level III - submicron CVT nanopores, shown in surface electron image. For illustration, the size of the micrographs (c) and (d) are illustrated as red boxes in (b) and (c), respectively. FIG.4. In situ XRD results for Fe-25W foam during the first two redox cycles. (a) waterfall diffractograms as a function of time for the full two-cycle process, with phases and reduction (“red”) and oxidation (“ox”) reactions indicated. (b) plot of evolution of normalized peak height integrals for Attorney Docket No.: 616146.100555 metallic phases during the first and second cycles (c) plot of evolution of normalized peak integrals for oxide phases during the first and second cycle. FIG.5. Radial cross-section SEM micrographs showing microstructural evolution for Fe-25W foams after 1, 10, 20, 50, and 100 redox cycles at 800 ºC, in the reduced state (left column) and the oxidized state (right column). Red arrows mark micropores, and green arrows mark W-containing phases (Fe2W for the reduced state, FeWO4 for the oxidized state). FIG.6. Surface SEM micrographs of Fe-25W lamellae after 1, 10, 20, 50, and 100 redox cycles in the reduced state (left column) and the oxidized state (right column). Red arrows mark micropores and green arrows mark W-rich features. FIG.7. XRD spectra taken from polished cross-sections of (a) reduced foams after 1, 10, 20, 50, and 100 cycles, with Fe and Fe2W peak positions indicated. The shaded gray region indicates the diffraction angle range where Fe2W peaks show nanocrystalline broadening, causing convolution with other peaks. (b) oxidized foams after 0, 10, 20, 50, and 100 cycles, with Fe3O4 and FeWO4 peak positions indicated. FIG.8. Axial cross-section mosaic optical macrographs showing evolution of lamellar architecture after 1, 10, 20, and 50 redox cycles. Blue dashed lines indicate the foam envelope volume, red circle indicates a region at the center of the foam where lamellar structure is no longer present. FIG.9. Axial cross-sectional SEM micrographs (left) and surface micrographs (right) of Fe- 25W tapped powder beds in the reduced state after (a, b) 1 cycle, (c, d) 10 cycles, (e, f) 20 cycles, (g, h) 50 cycles, and (i, j) 100 cycles. Green arrows point to relatively large initial Fe grain size, which shrinks to submicron (red arrows) with cycling. FIG.10. In situ XRD spectra for the first two redox cycles of Fe-25W powders, performed at 800 °C under flowing CO2 as oxidizing gas and Ar-4%H2 as reducing gas. (a) time-resolved XRD spectra with various phases identified during oxidation (ox) and reduction (red). (b) Normalized peak integrals for metallic phases: Fe (blue), µ-Fe7W6 (purple), α-W (cyan) and λ-Fe2W (black). (c) Normalized peak integrals for oxide phases: FeO (orange), FeWO4 (green), and Fe3O4 (red). FIG.11. (a-c) Scanning electron microscopy (SEM) images of the surface of a Fe powder bed (a) before oxidation, (b) after CO2 oxidation, in a region with both oxide (white dashed region) and metal (pink dashed region). (c) After CO2 oxidation, in a region with thoroughly sintered oxide surface. (d-f) SEM images of the surface of a Fe-25W powder bed (d) before oxidation, with a two-phase µ- Fe7W6+ α-Fe composition. (e) After CO2oxidation in a porous region (f) in a more densified region. Oxidation was performed for 4 h at 800 °C in the TGA instrument. FIG.12. TGA data (mass gain vs. time) for Fe-25W specimens subjected to CO2 oxidation upon Attorney Docket No.: 616146.100555 heating to, and holding at, 800 ºC (temperature profile is shown on the right axis) for their 1st, 10th, 100th, and 165thcycle, for specimens previously subjected to redox cycling outside the TGA. Also shown for comparison is the oxidation for a Fe powder bed for its 1stcycle. The grey and brown dashed 1 line indicate the theoretical mass gain for complete oxidation of Fe (Fe ^ 3 Fe3O4) and Fe-25W (Fe + 2 Fe2W ^ FeWO4 + 3 Fe3O4), respectively. The mass change is to the mass reading at t = 32 min; a control run with an empty TGA pan revealed an apparent mass gain (due to gas flow) up to t = 32 min (450 °C): thus, mass gain values measured before this point are inaccurate and are not shown. FIG.13. In situ XRD spectra for Fe-25W starting at cycle 11 at 900 ºC, under CO2 as oxidizing gas and under H2 as reducing gas. (a) time-resolved XRD spectra showing the phase transformations for 3 half cycles (oxidation-reduction-oxidation). (b) Normalized peak integrals for non-oxide phases: Fe (blue) and λ-Fe3W3C (purple). (c) Normalized peak integrals for oxide phases: FeO (orange), FeWO4 (green), Fe3O4 (red), and Fe2O3 (cyan). FIG.14. SEM images showing microstructural evolution of the surface of Fe-25W powder beds under CO2 / H2 redox cycling, after (a, b) 1 cycle, (c, d) 10 cycles, (e, f) 100 cycles, and (g, h) 165 cycles, in the (a, c, e, g) reduced and (b, d, f, h) oxidized states. Cyan arrows mark micron-size Fe regions in the first cycle, yellow arrows mark micropores, and red arrows mark submicron pores. Insets in (e-h) show submicron features at higher magnification. FIG.15. Cross-sectional SEM images of a Fe-25W powder bed portion comprising numerous powder agglomerates (100 µm-scale) separated by channels (~ 50 µm wide, yellow letter C). (a, d, g, j) macrostructure; (b, e, h, k) mesostructure; and (c, f, i, l) microstructure for various cycles: (a, b, c) uncycled, and after (d, e, f)10 cycles, (g, h, i) 100 cycles, and (j, k, l) 165 cycles. Blue arrows mark Fe regions, orange arrows mark Fe2W regions. FIG.16. (a) Fe-W phase diagram from CALPHAD-based calculation using Thermo-Cal software, showing compositions of the four Fe-W foams at their reduction temperature of 600 ºC (hollow symbols) and at their subsequent sintering temperatures (1000, 1100, and 1200 ºC, solid symbols). (b) X-ray diffraction patterns of as-sintered Fe-W foams, with phases identified. The reported diffraction patterns of Fe7W6and Fe2W phases are plotted as well. FIG.17. In situ XRD characterization of Fe-10W foam during the first redox cycling at 800 ⁰C. (a) Evolution of XRD patterns as a function of time for the first oxidation / reduction cycle. (b,c) Evolution of normalized peak intensity integrals during the first oxidation / reduction cycle for (b) metallic phases and (c) oxide phases. FIG.18. In situ XRD characterization of Fe-18W foam during the first redox cycling at 800 ⁰C. (a) Evolution of XRD patterns as a function of time for the first oxidation / reduction cycle. (b, c) Attorney Docket No.: 616146.100555 Evolution of normalized peak intensity integrals during the first oxidation / reduction cycle for (b) metallic phases and (c) oxide phases. FIG.19. In situ XRD characterization of Fe-33W foam during the first redox cycling at 800 ⁰C. (a) Evolution of XRD patterns as a function of time for the first oxidation / reduction cycle. (b, c) Evolution of normalized peak intensity integrals during the first oxidation / reduction cycle for (b) metallic phases and (c) oxide phases. FIG.20. Backscattered electron (BSE) micrographs of lamellar surface in the Fe-W foams: (a) Fe-10W, (b) Fe-18W, (c) Fe-25W, and (d) Fe-33W. The α-Fe, λ-Fe2W, and μ-Fe7W6 phases are marked in each foam, and micropores have black contrast. Submicron porosities in the cross-section view of Fe-18W foam are marked with white arrows in the higher-magnification insert in (b). FIG.21. BSE micrographs of the cross-section of lamellar structures of (a, b) Fe-10W, (c, d) Fe-18W and (e, f) Fe-25W foams in the reduced state after 20 redox cycles at 800 ⁰C. A ligament within a porous lamella - including α-Fe and λ-Fe2W phases and microporosity (black contrast) - is highlighted with a red dashed circle. FIG.22. BSE micrographs of the surface of Fe-10W lamellae in the oxidized state after (a, b) 1 redox cycle and (c, d) 20 redox cycles at 800 ⁰C. (a) Cross-section of an oxidized lamella after the first cycle; in this inclined view, the white dashed lines indicate the edges of the lamella. (b) A magnified view of the lamella cross-section in (a). (c) Cross-section of two oxidized lamellae separated by a channel, after 20 redox cycles. The yellow dashed lines show the edges of Fe3O4 shells formed at the lamellae surface. For lamella #2, two lamellae (#2a and #2b) contacted and sintered together on the top part but left a fork at the bottom marked with a green arrow. (d) An inclined view showing the side surface of an oxidized lamellae, with cracks in the shell providing gas access to the porous core. FIG.23. BSE micrographs of the cross-section of a Fe-18W porous lamella in the oxidized state after (a) 1 cycle and (b) 20 cycles at 800 ⁰C. FIG.24. BSE micrographs of the cross-section of lamellar structures of Fe-18W foams in the reduced state after (a, b) 1 cycle, (c, d) 25 cycles, (e, f) 30 cycles, and (g, h) 50 redox cycles, illustrating lamellar buckling, as well as increasing microporosity and decreasing ligament size within lamellae with increasing cycling numbers. Ligaments within a porous lamella - including α-Fe and λ-Fe2W phases and microporosity (black contrast) - are highlighted with red dashed circles. FIG.25. XRD patterns of Fe-18W foams in their reduced, metallic state, after 1, 25, 30, and 50 redox cycles. The hump marked with a grey shadow shows strong peak broadening of the nanocrystalline λ-Fe2W phase. FIG.26. BSE micrographs of the radial cross-section at the edge of Fe-10W foams, showing Attorney Docket No.: 616146.100555 shell formation in (a, b) the reduced state and (c, d) oxidized state after 20 redox cycles at 800 ºC. The edge (surface) of the foams is marked with a white dashed line. FIG.27. Stitched BSE micrographs of an axial cross-section of a full Fe-18W foams after 20 redox cycles showing lamellar buckling (e.g., within the yellow dotted frame and partial densification (white dotted ellipse). The yellow vertical arrow indicates the direction of gravity (g) during redox cycling. The blue horizontal arrows indicate the deflection of buckled lamellae in the radial direction due to brooming deformation. FIG.28. Schematic illustrating evolution processes during redox cycling for lamellar Fe-W foams with (b-d) low W and (e-f) high W content. The white region within the lamellae represents micro-porosity from freeze casting and partial sintering. Colored regions represent different phases: α- Fe in orange, λ-Fe2W in green, FeWO4in red, and Fe3O4in yellow. (a) 3D view of two lamellae in an as-sintered foam. A magnified view of the cross-section of low-W lamellae is shown in the following panels: (b) Initial microstructure showing uniform mixture of α-Fe and λ-Fe2W phases in as-sintered lamella with microchannels; (c) Core-shell structure, with dense Fe3O4shell and porous core (comprising Fe3O4 and FeWO4 phases) in the oxidized state; (d) Metallic core-shell structure with dense α-Fe shell and porous α-Fe+λ-Fe2W core in the reduced state. A magnified view of the cross- section of high-W lamellae is shown in the following panels: (e) Initial microstructure of as-sintered lamella showing microchannels and ligaments with a uniform mixture of α-Fe and λ-Fe2W phases; (f) ligaments showing an homogeneous distribution of Fe3O4 and FeWO4 phases; (g) refined, uniform mixture of α-Fe+λ-Fe2W phases in ligaments in the reduced state. FIG.29. Oxygen and hydrogen capacities, electron number, and theoretical Nerst potential (EN) as a function of W concentration at 800oC. The calculation is performed for a total of 1 mole of Fe plus W atoms. The left y-axis is the oxygen (O) and hydrogen (H) capacity calculated for 1 mole of Fe+W atoms, and the right y-axis is the corresponding capacity per alloy unit mass, plotted in (a) and (b). FIG.30. Backscattered electron (BSE) and secondary electron (SE) micrographs of an identical region on the surface of Fe-33W foam. FIG.31. BSE micrographs of radial cross-section of as-sintered foams, showing lamellar structure with epoxy-filled channels with black contrast: (a, b) Fe-10W and (c, d) Fe-18W. FIG.32. SEM micrographs of Fe-25W in the reduced state after 50 redox cycles at 800oC, showing (a, b) cross-section and (c, d) surface of foam. The lamellar structure of the freeze-cast foams evolves a into near-isotropic, highly-porous, powder network after 50 cycles. FIG.33. In situ X-Ray diffraction data for redox cycling of Fe-20Ni-20W lamellae at 800 °C. (a) Waterfall XRD spectra are shown as a function of time for the first two redox cycles, with all phases Attorney Docket No.: 616146.100555 labelled. (b) Normalized peak integrals for metallic phases: µ-Fe7W6(black), γ-Fe(Ni) (dark blue), and γ-Ni(Fe) (cyan). The two γ phases are distinguished by their peak positions, with the Ni-rich γ-Ni(Fe) peak shifted to higher 2θ values as compared to the Fe-rich γ-Fe(Ni) peak; for this reason, the γ-Ni(Fe) signal is only analyzed when at a 2θ value distinct from the γ-Fe(Ni). (c) Normalized peak integrals for oxide phases: FeWO4 (green) and Fe3O4 (red). Green arrows in (b,c) mark instances of transient re- oxidation during the reduction half-cycles. FIG.34. Photographs of as-sintered Fe-W foams with various W concentrations. FIG.35. Micrographs of Fe-20Ni-20W freeze-cast foams after reduction and sintering. (a) radial cross-section electron image showing a macroscopic view of lamellae and channels, (b) radial cross- section electron image of channels (letter C) and lamellae (yellow letter L) (c) higher magnification electron image of lamella interior (radial foam cross-section) with ligaments (red letter L) and micropores (letter M) showing a two-phase microstructure (µ-Fe7W6 (white) and γ-Fe(Ni,W) (dark gray)), and (d) electron image of the surface of an unmounted lamella, showing microporous network made possible by sintering inhibition, with ligaments exhibiting a two-phase microstructure, shown in inset. FIG.36. Electron micrographs showing surface structure of a single Fe-20Ni-20W lamella after 1, 20, and 50 cycles in the reduced (top row) and oxidized (bottom row) states. Orange arrows mark micropores, red arrows mark submicron pores, red letter L marks examples of ligaments, and yellow arrows mark FeWO4 whiskers. DETAILED DESCRIPTION OF THE INVENTION The invention will now be described more fully hereinafter with reference to the accompanying drawings, in which exemplary embodiments of the invention are shown. However, this invention 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 specification will be thorough and complete and fully convey the invention's scope to those skilled in the art. Like reference numerals refer to like elements throughout. The terms used in this specification generally have their ordinary meanings in the art, within the context of the invention, and in the specific context where each term is used. Certain terms used to describe the invention are discussed below, or elsewhere in the specification, to provide additional guidance to the practitioner regarding the description. For convenience, certain terms may be highlighted, for example using italics and / or quotation marks. The use of highlighting has no influence on the scope and meaning of a term; the scope and meaning of a term are the same, in the same context, Attorney Docket No.: 616146.100555 whether or not it is highlighted. It will be appreciated that same thing can be said in more than one way. Consequently, alternative language and synonyms may be used for any one or more of the terms discussed herein, nor is any special significance to be placed upon whether or not a term is elaborated or discussed herein. Synonyms for certain terms are provided. A recital of one or more synonyms does not exclude the use of other synonyms. The use of examples anywhere in this specification including examples of any terms discussed herein is illustrative only, and in no way limits the scope and meaning of the invention or of any exemplified term. Likewise, the invention is not limited to various embodiments given in this specification. It will be understood that, as used in the description herein and throughout the claims that follow, the meaning of “a”, “an”, and “the” includes plural reference unless the context clearly dictates otherwise. Also, it will be understood that when an element is referred to as being “on” another element, it can be directly on the other element or intervening elements may be present therebetween. In contrast, when an element is referred to as being “directly on” another element, there are no intervening elements present. As used herein, the term “and / or” includes any and all combinations of one or more of the associated listed items. It will be understood that, although the terms first, second, third, etc. may be used herein to describe various elements, components, regions, layers and / or sections, these elements, components, regions, layers and / or sections should not be limited by these terms. These terms are only used to distinguish one element, component, region, layer or section from another element, component, region, layer or section. Thus, a first element, component, region, layer or section discussed below could be termed a second element, component, region, or section without departing from the invention's teachings. Furthermore, relative terms, such as “lower” or “bottom” and “upper” or “top,” may be used herein to describe one element’s relationship to another element as illustrated in the figures. It will be understood that relative terms are intended to encompass different orientations of the device in addition to the orientation depicted in the figures. For example, if the device in one of the figures. is turned over, elements described as being on the “lower” side of other elements would then be oriented on “upper” sides of the other elements. The exemplary term “lower”, can, therefore, encompasses both an orientation of “lower” and “upper,” depending on the particular orientation of the figure. Similarly, if the device in one of the figures is turned over, elements described as “below” or “beneath” other elements would then be oriented “above” the other elements. Therefore, the exemplary terms “below” or “beneath” can encompass both an orientation of above and below. It will be further understood that the terms “comprises” and / or “comprising,” or “includes” Attorney Docket No.: 616146.100555 and / or “including” or “has” and / or “having”, or “carry” and / or “carrying,” or “contain” and / or “containing,” or “involve” and / or “involving, and the like are to be open-ended, i.e., to mean including but not limited to. When used in this specification, they specify the presence of stated features, regions, integers, steps, operations, elements, and / or components, but do not preclude the presence or addition of one or more other features, regions, integers, steps, operations, elements, components, and / or groups thereof. Unless otherwise defined, all terms (including technical and scientific terms) used herein have the same meaning as commonly understood by one of ordinary skill in the art to which this invention belongs. It will be further understood that terms, such as those defined in commonly used dictionaries, should be interpreted as having a meaning that is consistent with their meaning in the context of the relevant art and this specification, and will not be interpreted in an idealized or overly formal sense unless expressly so defined herein. As used in this specification, “around”, “about”, “approximately” or “substantially” shall generally mean within 20 percent, preferably within 10 percent, and more preferably within 5 percent of a given value or range. Numerical quantities given herein are approximate, meaning that the term “around”, “about”, “approximately” or “substantially” can be inferred if not expressly stated. As used in this specification, the phrase “at least one of A, B, and C” should be construed to mean a logical (A or B or C), using a non-exclusive logical OR. As used herein, the term “and / or” includes any and all combinations of one or more of the associated listed items. It should be noted that the term “ligament”, used in the disclosure, refers to the solid content of the bed (as opposed to pores) after the powders sinter together. The description below is merely illustrative in nature and is in no way intended to limit the invention, its application, or uses. The broad teachings of the invention can be implemented in a variety of forms. Therefore, while this invention includes particular examples, the true scope of the invention should not be so limited since other modifications will become apparent upon a study of the drawings, the specification, and the following claims. For purposes of clarity, the same reference numbers will be used in the drawings to identify similar elements. It should be understood that one or more steps within a method may be executed in a different order (or concurrently) without altering the principles of the invention. Grid-scale implementation of renewable energy production relying on intermittent sources, (e.g., solar and wind) requires an energy storage system that allows excess electrical energy produced during high-production times to be stored and released later at low-production times, since the natural power source may not align with the demand for energy. This energy storage material should be Attorney Docket No.: 616146.100555 inexpensive and produced at scale; iron-air solid-oxide flow batteries fill this need since iron is extremely abundant and inexpensive as compared to other battery materials (e.g., Li, Co and Ni), however, most current iron-based systems are hindered by short cycle lifetime. In addition, the rechargeable oxide battery (ROB), based on the high-temperature redox of an Fe-based material to reversibly fuel a high temperature fuel cell is one promising nascent technology that has been hindered by fast degradation due to high temperature sintering, leading to a short lifespan. To address the degradation, this invention discloses an iron-tungsten redox system including an extremely resilient Fe-xW (at%) powder bed that can maintain the same structure and reaction kinetics for many more cycles than other Fe-based systems. The Fe-xW powder bed shows high resistance to degradation due to sintering during high temperature redox cycling, as listed in Table 1. This resistance comes from both the sintering inhibition inherent to W, and the nature of the intermetallic and mixed oxide compounds formed between Fe and W, such that W does not segregate and remains atomically mixed with Fe at all times, unlike previously studied Fe-Mo foams. It is the sintering inhibition and redox activity inherent to W, combined with a unique chemical vapor transport reduction mechanism of the mixed oxide FeWO4 that allows the powder bed to form and maintain a hierarchically porous structure that limits Fe sintering and maintains good reaction kinetics. This bed can serve as an excellent energy storage material for the ROB, while also being applicable to the related technologies of chemical looping combustion and high temperature CO2 utilization, both of which operate on the same high temperature redox reactions as the ROB. Table 1: Porosity and reactivity of different Fe-X porous bodies Composition 0 cycle 10 cycle 50 cycle This porosity (%) porosity (%) reactivity (%) Invention Fe 64 17 N / A NFe-25Ni 69 30 N / A NFe-25Co 62 46 N / A NFe-30Cu 65 14 N / A NFe-25Mo 79 73 45 NFe-25W 76 78 100 YFe-10W 74 42 N / A YFe-18W 75 70 77 YFe-33W 80 75 N / A YFe-20Ni-20W 75 77 55 YFe-9Mo-9W 70 37 80 YTable 1 displays porosity and reactivity data for relevant Fe porous body compositions that have been tested. Pure Fe and the binary systems Fe-25Ni, Fe-25Co, and Fe-30Cu all show a rapid drop in porosity after 10 redox cycles, indicating poor performance. Fe-25Mo, Fe(10-33)W, Fe-20Ni-20W, Attorney Docket No.: 616146.100555 and Fe-9Mo-9W all show porosity retention or even increase after 10 cycles; these compositions are tested for reactivity after 50 cycles. Reactivity here refers to the fraction of oxide that can be reduced in 90 minutes exposure to flowing H2 at 800 °C. Fe-25W shows excellent porosity retention after 10 cycles, and excellent reactivity after 50 cycles. Fe-18W and Fe-9Mo-9W show similar porosity retention after 10 cycles, but slightly lower reactivity after 50 cycles. Fe-20Ni-20W similarly shows excellent porosity retention after 10 cycles, but lower reactivity. Fe-10W and Fe-33W show porosity retention after 10 cycles, btu have not been tested for reactivity after 50 cycles, In some embodiments, a tapped powder bed including Fe-25W is prepared using micron-scale oxide powders of Fe2O3 and WO3, which are reduced and sintered at high temperature into a metallic powder bed. Specifically, a powder mixture of Fe2O3and WO3precursors is produced by preparing a slurry of the Fe2O3and WO3powders in ethanol, which is then ball milled for 24 hours to mix the powder s and then dried. The dried powders are reduced at high temperature to produce a partially sintered metallic bed. During redox cycling at 800 C the bed is reversibly oxidized and reduced with no densification due to sintering, which would block gas access and slow reaction rates. The longevity of the Fe-25W beds is enabled by the sintering inhibition of W and the cyclic formation of submicron pores due to chemical vapor transport reduction of FeWO4. During redox cycling under steam (for oxidation) followed by hydrogen (for reduction), the bed displays hierarchical porosity: microscale pores enabled by high sintering resistance of W, and submicron pores enabled by chemical vapor reduction of the oxide phase FeWO4. These two levels of porosity are sufficient to prevent Fe and Fe3O4sintering in the bed, allowing for similar kinetics even after many full redox cycles, with the structure and stability intact after 100 cycles. Additionally, the oxidation of Fe and W into FeWO4and Fe3O4 lends this system a high rate of H2O-to-H2 conversion on a per-mole basis, higher than the pure Fe system or any system with inactive additive materials. This system is applicable to at least three distinct technologies, all of which utilize high temperature reduction and oxidation: i) rechargeable oxide battery, in which the Fe-25W converts steam into hydrogen to feed a fuel cell creating electric power, ii) chemical looping combustion reactor, in which the Fe-25W oxide reacts with methane (or other C-bearing fuels) to produce heat and a pure CO2waste stream, and iii) CO2utilization, in which the Fe-25W converts CO2 to CO for use as a chemical feedstock. Without intent to limit the scope of the invention, exemplary embodiments of the invention are given below. In one aspect of the invention, the redox system comprises a Fe-xW powder architecture fabricated to possess excellent resistance to degradation at a high-temperature during redox cycling between oxidized and reduced states, wherein x represents a concentration (at%) of tungsten (W) in the Attorney Docket No.: 616146.100555 powder architecture and is in a range of about 1-75 at%, wherein the powder architecture comprises powder bed, powder suspension, foams, fibers, and / or printed microlattice. In one embodiment, the redox cycling is performed with an oxidizing gas and an reducing gas alternatively, wherein the oxidizing gas are those that increase the oxidation state of iron (Fe) and / or tungsten (W) and comprises H2O, CO2, and / or O2, and wherein the reducing gas are those that decrease the oxidation state of Fe and / or W and comprises H2, hydrocarbons including CH4, ethane, propane and / or butane, NH3, and / or CO. In one embodiment, the oxidizing gas and the reducing gas are combined with any inert gases. In one embodiment, the high-temperature is in a range of about 500 – 1500 °C. In one embodiment, the powder architecture possesses the excellent resistance to degradation at about 800 °C during the redox cycling between the metallic and oxide states, with good reaction kinetics maintained up to a minimum of 100 full cycles. In one embodiment, the powder architecture is of Fe-1W, Fe-5W, Fe-10W, Fe-18W, Fe-20W, Fe-25W, Fe-30W, Fe-33W, Fe-40W, Fe-45W, Fe-50W, or Fe-75W. In one embodiment, the Fe-10W powder architecture with a lower volume fraction of W- containing phases of λ-Fe2W and FeWO4 has lamellae densification and core-shell structure formation, due to Fe outward diffusion during oxidation. In one embodiment, each of the Fe-18W and Fe-25W powder architectures comprises a sufficient volume fraction of W-rich phases including λ-Fe2W to inhibit sintering for α-Fe in the reduced state and FeWO4to inhibit sintering for Fe3O4in the oxidized state, thereby forming ligaments comprising two phases of Fe / λ-Fe2W and Fe3O4 / FeWO4, respectively. In one embodiment, each ligament in the Fe-25W powder architecture initially includes a mixture of bcc α-Fe(W) and µ-Fe7W6; during oxidation, α-Fe oxidizes to Fe3O4 and µ-Fe7W6 oxidizes to FeWO4and minor amounts of Fe3O4; and during reduction, the Fe3O4reduces back to bcc Fe, and the FeWO4 reduces more slowly to Fe2W. In one embodiment, there a mismatch in the stoichiometric ratio between Fe and W in moving from FeWO4(Fe:W=1) to Fe2W (Fe:W=2), whereby Fe from the surrounding matrix from prior reduction of Fe3O4 is used to form the intermetallic Fe2W. In one embodiment, the second oxidation is, on average, about 40-80% faster than the first oxidation, and the second reduction is, on average, about 10-50% faster than the first reduction. In one embodiment, in the reduced state, the powder architecture includes powders which exhibit 2-10 µm ligaments interpenetrated by open porosity of the same size, wherein each ligament, in turn, includes an α-Fe(-1.6W) backbone decorated, and at times interrupted, with submicron (1-2000 Attorney Docket No.: 616146.100555 nm) λ-Fe2W grains and submicron CVT pores, and wherein this same reduced microstructure is maintained through 100 cycles with the micron-scale open porosity increasing with cycle number. In one embodiment, the oxidized state includes two interpenetrating networks of Fe3O4 and FeWO4. In one embodiment, the powder architecture, after the first cycle, comprises three levels of pore sizes including macroscopic freeze-cast channels (level I) separating lamellae which are distributed into colonies with aligned radial orientation; microscale sintering-inhibition pore network (level II) within each lamella, fully continuous, formed due to the sintering inhibition of W; and submicron chemical vapor transport (CVT) pores (level III), formed within each W-containing area in the lamellae, regenerated at each reduction half-cycle, thereby creating a hierarchical porosity structure in the powder architecture. In one embodiment, the microscale sintering-inhibition pores increase the surface area and thus reactivity of the ligaments and the submicron CVT pores greatly increase surface area and limit damage accumulation by effectively resetting the microstructure after each cycle, producing nanocrystalline λ– Fe2W with corresponding nanopores. In one embodiment, the microstructural stability allows for the powder architecture to grow progressively more porous (rather than denser) with cycling, and the W-rich phases become both smaller in morphology and more homogeneously distributed throughout the ligaments, thereby increasing the sintering inhibition effect and generating more submicron pores to accelerate subsequent reaction. In one embodiment, the resistance is attributed to the sintering inhibition of W combined with the CVT mechanism of reduction. In one embodiment, the resistance is attributed to both the sintering inhibition inherent to W, and the nature of the intermetallic and mixed oxide compounds formed between Fe and W, such that W does not segregate and remains atomically mixed with Fe at all times. In one embodiment, W is atomically mixed with Fe, either in the intermetallic λ–Fe2W or in the mixed oxide FeWO4phases, thereby effectively preventing W from segregating into a pure W phase that reduces the degradation resistance of the Fe-25W powder architecture. In one embodiment, increasing W concentration has a minor effect on the oxidation rate but significantly extends the initial reduction time in the first cycle due to the slower reduction of more stable FeWO4, as compared to Fe3O4. In one embodiment, the chemical vapor transport (CVT) mechanism during reduction refines stable W-rich phases to a nanocrystalline scale, boosting the reaction kinetics of subsequent redox Attorney Docket No.: 616146.100555 cycles. The reduction time could be adjusted by designing W concentrations and associated phase fractions of FeWO4. In one embodiment, higher W concentration enhances the degradation resistance of Fe-W lamellar foams during high-temperature redox cycling, resulting in a more stable lamellae structure. In one embodiment, the powder architecture further comprises Ni, Co, Cu, and / or Mo, in compositional ranges Fe-(5-33)(Mo+W), Fe-(5-33)(Ni+W), or Fe-(5-33)(Mo+Ni+W), all in at%. In one embodiment, the powder architecture is of Fe-(5-33)(Mo, W), wherein Mo and W are both active in oxidation and reduction. In one embodiment, the powder architecture is of Fe-9Mo-9W. In one embodiment, the powder architecture is of Fe-(5-33)(Ni, W), wherein W is active in oxidation and reduction while Ni is inert. In one embodiment, the powder architecture is of Fe-20Ni-20W. In one embodiment, the Fe-20Ni-20W powder architecture, after initial reduction, include a two-phase mixture of µ-Fe7W6and γ-Fe(Ni,W), with significant microporosity due to the sintering inhibition of W, and wherein during oxidation at 800 °C, these phases are oxidized to a three-phase mixture of (Fe,Ni)WO4, Fe3O4, and γ-Ni(Fe); upon subsequent reduction by H2, the powder architecture returns to their initial composition. In another aspect, the invention relates to a device comprising the redox system as disclosed above. In a further aspect, the invention relates to a method for fabricating a redox system comprising milling a blend of an iron (Fe) or Fe-containing powder and a tungsten (W) or W-containing powder into a homogenous powder with Fe-xW composition; and, optionally, heat treating the powder blend to yield a powder architecture of Fe-xW, wherein x represents a concentration (at%) of W in the powder architecture and is in a range of about 1-75 at%. In one embodiment, the Fe or Fe-containing powder comprises a Fe powder, a Fe-rich powder, a Fe-containing oxide powder, a Fe-containing nitride powder, and / or a Fe-containing carbide powder. The W or W-containing powder comprises a W powder, a W-rich powder, a W-containing oxide powder, a W-containing nitride powder, a W-containing carbide powder, and / or ammonium paratungstate ((NH4)10(H2W12O42)·4H2O)). In one embodiment, the W or W-containing powder is comprised of any of W-rich chemical compounds, such as oxide, hydroxide, nitride, carbide and other organic variants, e.g., tungsten trioxide(WO3), tungstic acid (H₂WO₄), potassium tungstate (K₂WO₄), sodium tungstate (Na₂WO₄), and / or morecomplex compounds such as ammonium paratungstate. Attorney Docket No.: 616146.100555 In one embodiment, the powder architecture is of Fe-1W, Fe-5W, Fe-10W, Fe-18W, Fe-20W, Fe-25W, Fe-30W, Fe-33W, Fe-40W, Fe-45W, Fe-50W, or Fe-75W. In one embodiment, the powder architecture possesses the excellent resistance to degradation at about 800 °C during redox cycling between oxidized and reduced states, with good reaction kinetics maintained up to a minimum of 100 full cycles. In one aspect, the invention relates to a method for fabricating a redox system comprising providing a Fe-xW alloy; atomizing the alloy to form a powder; and optionally heat treating the powder to yield a powder architecture of Fe-xW, wherein x represents a concentration (at%) of W in the powder architecture and is in a range of about 1-75 at%. In one embodiment, said providing the Fe-xW alloy is performed by melting iron and ferrotungsten. In another aspect, the invention relates to a method for fabricating a redox system comprising producing a Fe-and W-rich oxide by sintering of precursor powders; pulverizing the Fe-and W-rich oxide to form a powder with Fe-xW oxide composition; and optionally heat treating the powder to yield a powder architecture with metallic Fe-xW composition, wherein x represents a concentration (at%) of W in the powder architecture and is in a range of about 1-75 at%. The invention, among other things, provides the following advantages: inexpensive, non toxic, earth-abundant materials (iron oxide and tungsten oxide); low-cost, clean processing (water carrier and sintering); hierarchical porous structure to maintain good reaction kinetics; no agglomeration of the alloying element during redox cycling; and no architecture required, decreasing processing cost and complexity. The invention may have widespread applications in solid-oxide iron-air flow batteries (reversible, multi cycles), chemical looping combustion reactor, CO2 utilization (convert CO2 to CO), hydrogen generator (water splitting), hydrogen storage, and the like. These and other aspects of the invention are further described below. Without intent to limit the scope of the invention, exemplary instruments, apparatus, methods, and their related results according to the embodiments of the invention are given below. Note that titles or subtitles may be used in the examples for convenience of a reader, which in no way should limit the scope of the invention. Moreover, certain theories are proposed and disclosed herein; however, in no way they, whether they are right or wrong, should limit the scope of the invention so long as the invention is practiced according to the invention without regard for any particular theory or scheme of action. EXAMPLE 1: Attorney Docket No.: 616146.100555 TUNGSTEN STRONGLY INHIBITS SINTERING OF POROUS IRON DURING HIGH- TEMPERATURE REDOX CYCLING This example discloses freeze-cast Fe-25W (at%) lamellar foams showing excellent resistance to degradation at 800 °C during steam-hydrogen redox cycling between the metallic and oxide states, with fast reaction kinetics maintained up to at least 100 redox cycles with full Fe utilization. This very high stability stems from the sintering inhibition of W combined with the freeze-cast architecture and the chemical vapor transport (CVT) mechanism of reduction. These three factors create a hierarchical porosity in the foam, including (i) macroscopic elongated channels, (ii) micro-scale sintering inhibition pores, and (iii) submicron CVT pores. Microstructural characterization via SEM and EDS is combined with in situ XRD to fully explore the phase evolution and microstructural impact of W on Fe during redox cycling. Comparison with tapped Fe-25W (at%) powder beds reveals that the freeze-cast channels and lamellae are not critical to the performance of the material. In the example, the composition is chosen to form continuous phases of Fe2W and FeWO4 in the material, allowing for detailed examination of the effects of the two-phase microstructure as shown in our previous studies with the same content of Ni, Co, and Mo. This relatively high W content also serves as a useful point of comparison for future studies of lower W content. The degradation resistance of this material stems from both the sintering inhibition inherent to W and the nature of the intermetallic and mixed oxide compounds formed between Fe and W, such that W does not segregate, and remains atomically mixed with Fe at all times, unlike previously studied Fe-Mo foams. The Fe-W foams show good redox kinetics over long-term cycling: they fully oxidize and reduce within 90 min after 100 full cycles, with no segregation of W or densification of the foam observed. Methods Foam fabrication: Fe-25W (at%) freeze-cast foams were prepared by aqueous freeze casting, following procedures developed in previous work for Fe-Ni, -Co, -Cu and -Mo foams. An aqueous slurry was prepared, containing 5.84 vol% Fe2O3 (Noah Technologies, 99.9%, < 3µm), 4.15 vol% WO3 (SSNano, 99.5%, < 100 nm) and 87.46 vol% DI water. By mass, the Fe2O3 / WO3ratio is 0.965. Additionally, 0.55 vol% Zephrym PD 4974 (Croda) was added as a dispersant. This slurry was ball- milled for 24 h with yttria-stabilized zirconia milling media on a rotary ball mill. After milling, 2.00 vol% polyethylene glycol was added as a binder. The prepared slurry was degassed and then chilled to 0 °C in an ice bath before freeze-casting. Directional freeze casting was performed on a copper plate chilled by a thermoelectric cooler (Mauser). The slurry was poured into a Teflon mold (25 mm OD, 15 mm ID, 15 mm height) and the Attorney Docket No.: 616146.100555 temperature of the copper base of the plate was reduced to -30 °C following an exponential cooling curve; this cooling profile yields a constant rate of ice growth such that the formed lamellae maintain a constant thickness and spacing throughout the height of the foam. After freeze casting, the lowest 1 mm of the frozen foam was removed by razor blade. The cast foams were freeze-dried (0.13 mbar residual pressure) for 24 h at -54 °C to remove all ice before further processing. After freeze drying, the oxide green bodies were heat treated to yield metallic foams. First, the organic material was burnt out at 300 °C for 1 h, then the oxides were reduced with H2at 600 °C for 4 h, and finally the foams were sintered at 1200 °C for 3 h (10 °C / min heating and 5 °C / min cooling rates throughout) under flowing H2 (UHP, Airgas). The average foam mass after reduction and sintering was 0.83 g. To prepare tapped powder beds for comparison, the slurry was prepared as described above for the freeze-cast foams; it was then dried in laboratory air for 72 h rather than freeze cast, and then ground to powder by hand with a mortar and pestle. The same heat treatment as for the freeze-cast foams was used on these powders, which were kept in an alumina tube (9 mm ID) to maintain a cylindrical shape while in the furnace. Foam redox cycling: Redox cycling was carried out in an alumina tube furnace (20 mm ID, 500 mm length) at 800 °C (10 °C / min heating rate) fed by both H2and H2O. Pure H2(UHP, Airgas) was delivered directly by a mass flow controller at 200 sccm. Steam (H2O) was delivered by flowing Ar (UHP, Airgas) as a carrier gas at 120 sccm through bubbling in water heated to 93 °C (to provide an H2O pressure of 770 mbar, sufficient to oxidize Fe to Fe3O4). Alternating cycles of H2and H2O exposure, 90 min each, were controlled by a LabVIEW®program. The gas used during cooldown was either 100 sccm H2 (for foams in the reduced state) or 20 sccm Ar-4% H2 (for foams in the oxidized state). A cooling rate of 5 °C / min was used, and the foams were brought to room temperature before removing from the furnace. The freeze cast foams were unconstrained, and thus free to expand and contract during the redox cycles. For the powder beds, two sample sets were fabricated: one set was taken out of the alumina tube (9 mm diameter) and was free to expand and contract like the freeze cast foams, the other set was left in the alumina tube throughout cycling, thus being radially constrained. Foam microstructural characterization: To obtain cross-sections, foams were mounted in epoxy and vacuum infiltrated to fill pores before grinding and polishing to 1 µm. The mounted foams were then ion milled with a Leica TIC3X to improve surface visibility. Polished samples were sputter- coated with 9 nm Au / Pd for SEM characterization. SEM characterization was performed on Hitachi SU8030 on either unmounted foams (with as-fabricated surfaces) or mounted and polished foam cross- sections. Attorney Docket No.: 616146.100555 X-ray characterization: In-situ X-ray diffraction spectra recorded during redox cycling of freeze-cast specimens were collected in a similar setup to our previous work, with minor modifications. Spectra were collected on a Stadi-MP (Stoe, Germany) instrument, with an asymmetric curved Ge monochromator under pure Ag-Ka1radiation (λ = 0.56 Å) and a one-dimensional silicon strip-detector (MYTHEN2 1k, from Dectris, Switzerland), operated at 40 kV and 40 mA (Beam Size 4 x 0.8 mm). Data were collected in Debye-Scherrer (transmission) geometry, using 1-minute scans with 2θ diffraction angles spanning 5.7 to 24.3º, after calibration against a NIST Si standard (640d). A colony of metallic lamellae was extracted, using a razor blade, from a reduced and sintered Fe-25W foam, measuring 1-1.5 mm in height and 1-1.5 mm in thickness. The lamellae were introduced in a 1.5 mm diameter quartz capillary and surrounded on both ends by amorphous quartz wool. The remainder of the capillary volume on the outlet end of the sample was filled with porous ceramic blocks to prevent sample movement upon changes in gas flow. The capillary was installed into a water-cooled, graphite- heated furnace, with temperature stability of 0.1 °C. The experiment began by heating the sample to 800°C (30°C / min) under flowing Ar-4% H2, to ensure the sample remained metallic until the cycling temperature was reached. Once at 800°C, the gas was switched to Ar, bubbled through a 40 °C water bubbler (PH2O = 73 mbar) for oxidation. All gas lines between the bubbler and capillary were heated to >40 °C to prevent steam condensation. Following complete oxidation, noted by unchanging diffraction patterns, Ar-4% H2 was flowed for reduction. This diluted reduction gas was used to slow the reduction reaction to match the time- resolution of the data collection. This procedure was repeated for two consecutive oxidation-reduction cycles. Processing of diffraction patterns was done in Python, with waterfall plots of diffraction spectra, after background correction with a modified polynomial fit. For each phase of interest, the strongest diffraction peak, indexed using reference patterns from the Inorganic Crystal Structure Database (ICSD), was fitted to a Pseudo-Gaussian model using the Lmfit package. These fitted peaks were then integrated to achieve an area, assumed to be proportional to the volume of the diffracting phase, and then normalized to the respective maxima of each phase. Ex-situ X-ray diffraction spectra were gathered in reflection mode from polished cross-sections on a Rigaku Smartlab Gen 2 with a Cu source. Data processing was performed in OriginLab to remove the epoxy background signal from the XRD spectra. The thermodynamic analysis of the Fe-25W system was performed using the energy and enthalpy of formation values experimentally measured for Fe2W and FeWO4, as well as standard values for Fe3O4 and H2O from the NIST-JANAF thermochemical tables, following a standard procedure. Attorney Docket No.: 616146.100555 Results and Discussion Thermodynamic Considerations: The isothermal (800 °C) section of the Fe-W-O ternary phase diagram is shown in FIG.1. At this temperature, the fully reduced Fe-25W foam is expected to include an equimolar mixture of two phases: bcc α–Fe(1.3W) and λ–Fe2W. During oxidation by H2O, this mixture approaches the Fe3O4 – FeWO4 two-phase region in the ternary phase diagram (FIG.1, (a), blue arrow), first moving into a three-phase Fe-Fe2W-FeWO4 region, followed by a FeO – Fe3O4 – FeWO4region, and finally the two-phase Fe3O4– FeWO4region (FIG.1, (a), green line). Further oxidation with steam to form Fe2O3 is not possible under ambient pressure conditions. During H2 reduction, oxygen is removed from the system as H2O, and the phase evolution proceeds in reverse, returning to the initial equimolar Fe+Fe2W composition. While the ternary phase diagram describes oxygen additions to the Fe-25W composition accurately, the use of steam as an oxidizing agent and hydrogen as a reduction agent brings additional thermodynamic considerations. The Baur-Glässner diagram is a useful tool showing phase stability as a function of H2 / H2O ratio and temperature. Oxidation reactions that require pure (or nearly pure) H2O, and reduction reactions that require pure (or nearly pure) H2 are not practical for use in rechargeable oxide batteries, since the energy storage material is expected to be exposed to a mixture of both gases. FIG.1, (b) shows the Baur-Glässner diagram for the Fe system (black, blue, and green lines), with the 2 Fe-25W system (purple line for Fe2W + Fe → FeWO4 + 3Fe3O4) superposed, based on experimentally- determined thermodynamic energies. The diagram indicates that, while FeWO4is more stable than Fe3O4, it is not so stable that pure H2 is necessary for its reduction at 800 ºC (in fact, a H2 / H2O ratio of 2.8 is sufficient at 800 ºC). This is corroborated by the conversion of the Baur-Glaessner diagram to equivalent oxygen partial pressure, shown in FIG.1, (c). A suitable energy-storage material should have its lower and upper oxidation state equilibrium lines lie between those for a 4:1 and 1:4 mixture of H2 / H2O (FIG.1, (c), red dashed lines), taken as values that can be easily achieved in battery operation. The oxygen partial pressure pO2 values for Fe2W / FeWO4 (purple line) closely approach those for Fe3O4 / FeO (green line) at 1073 K, indicating that the W addition does not reduce the suitability of the Fe redox system from a thermodynamic perspective. Using the Nernst equation for high-temperature fuel cells, these pO2 values give a Nernst potential of 0.92 V for the Fe-25W alloy, slightly lower than the 0.94 V value determined for the pure Fe system. Initial structure: To provide sufficient mechanical strength to the foam so they are not damaged during handling, sintering is performed at a much higher temperature (1200 °C) than redox cycling (800 °C). As a result, the initial phase composition is different than that achieved during cycling. After reduction and sintering at 1200 °C, each lamella in the foam includes a mixture of bcc a-Fe (with ~1 Attorney Docket No.: 616146.100555 at% in solid solution) and µ–Fe7W6, as predicted by the Fe-W phase diagram. During oxidation (the first half cycle), a-Fe oxidizes to Fe3O4 and µ-Fe7W6 oxidizes to FeWO4 and minor amounts of Fe3O4. The subsequent reduction of these two oxide phases at 800 ºC does not return the foam to its initial phases (reached at 1200 ºC); rather, the 800 ºC equilibrium phases are reached: bcc a-Fe (with 1.3W in solid solution) and λ-Fe2W. During reduction, the Fe3O4 reduces back to bcc Fe, and the FeWO4 reduces more slowly to Fe2W. Because there is a mismatch in the stoichiometric ratio between Fe and W in moving from FeWO4(Fe:W=1) to Fe2W (Fe:W=2), Fe from the surrounding matrix (from prior reduction of Fe3O4) is used to form the intermetallic Fe2W. Reduction of Fe3O4 powders is well described by the shrinking-core model, where the reducing gas (H2) dissociates and absorbs into vacancies in the oxide surface, before reacting with oxygen anions supplied from Fe3O4to form H2O and Fe cations, with charge balance maintained by the donation of electrons from the oxygen anion. The H2O molecule then desorbs while the Fe ions and electrons migrate to the internal Fe3O4 / FeO interface to form FeO. While the first Fe3O4→FeO reduction is dependent on the inward diffusion of Fe ions, the subsequent FeO→Fe reduction relies on the outward diffusion of oxygen anions through the outer metallic Fe layer to the gas interface; as the Fe layer grows, the diffusion distance increases, thus slowing the reaction kinetics. Short-circuit diffusion via cracks, pores, or grain boundaries can however improve kinetics. In the Fe-25W foams studied here, the high internal porosity of each lamella provides unobstructed gas access to each micron-scale Fe3O4 region, with diffusion distances being greatly shortened as compared to larger, dense Fe particles. The mechanism of FeWO4reduction, however, is different from that of Fe3O4: at high temperatures and in the presence of H2, FeWO4may be reduced by a chemical vapor transport (CVT) mechanism, as also observed for hydrogen reduction of WO3 and WO2. The observation of CVT for WOx reduction implies that the mixed oxide FeWO4 must undergo a transient state where the WOx content drops because of CVT, and the FeO4-xcontent decreases via the shrinking core mechanism. Despite this, no pure W phase is observed during reduction (as observed via XRD and EDS), only the equilibrium λ-Fe2W phase. Thus, it appears that, while some pure W may initially form due to the CVT reduction of binary WOx, it reacts with nearby Fe to form the binary λ-Fe2W phase. The absence of metallic W from XRD and EDS measurements does not preclude that small amounts of transient W form, below the detection limit, and react to λ-Fe2W. The main consequence of the CVT reduction hypothesis is that reduced λ-Fe2W will be nanocrystalline, with submicron pores present between the nascent Fe2W nano-grains. The CVT mechanism is consistent with the microstructure of the lamellae before (FIG.2, (a)) and after (FIG.2, (b)) the first reduction half cycle at 800 ºC, and the corresponding XRD patterns shown in FIG.2. Red Attorney Docket No.: 616146.100555 arrows mark α-Fe (with some W in solid solution), blue arrows mark μ-Fe7W6before the first cycle, and yellow arrows mark λ-Fe2W after the first cycle, as determined by backscatter electron Z-contrast in conjunction with XRD patterns. The XRD patterns (FIG.2, (c)) confirm that the λ-Fe2W phase is nanocrystalline (as shown by the broad shoulder), as well as highly textured, as the peak at diffraction angle 2q = 38° is relatively narrow compared to the broad nanocrystalline peaks between 40 and 48°. Additionally, a small amount of ternary carbide (Fe3W3C) is initially present, formed from reaction between the organic binder and the foam during burnout, but this carbon is removed by oxidation to CO / CO2 during the first oxidation half-cycle. This nano-porous and nanograin microstructure provides a further boost to the kinetics of subsequent redox cycles, as the slower-reacting W-rich phase is nanometric, increasing its kinetics of oxidation and reduction. The submicron pores also provide a greater degree of gas access to the foam, thus lowering solid-state diffusion distances to submicron levels. After the foam phases equilibrates in the first cycle, subsequent redox cycling is expected to follow the following reversible reaction: ^^^^ + ^^^^20 2 20 2^^ + ^^2^^ ↔ ^^^^^^^^4 +^^^^3^^4+ ^^2proceeding forward A crucial consequence of this reaction sequence is that W is always found in a compound with Fe, and no Fe-free W or WOx compounds are expected to form. After the first cycle, the foam includes three levels of pore size, as shown schematically in FIG. 3, (a), and in micrographs in FIG.3, (b)-(d): (level I) macroscopic freeze-cast channels, ~20 μm in width and running the full height (12 mm) of the foam, separating lamellae which are distributed into colonies with aligned radial orientation; (level II) microscale sintering-inhibition pore network within each lamella, fully continuous, formed due to the sintering inhibition of W; (level III) submicron CVT pores, formed within each W-containing area in the lamellae, regenerated at each reduction half-cycle. The retention of these three levels of porosity is key for the prolonged cycling stability of the foam, as this pore hierarchy provides excellent gas access to all parts of the foam, minimizing the reaction time needed to fully oxidize or reduce it. Phase evolution of freeze-cast foams during redox cycling: To confirm the expected phase evolution and compare the kinetics of oxidation and reduction of these various phases, in situ x-ray diffraction was performed on lamellae cut from the full freeze cast foams; to slow the reaction kinetics for better time resolution, flow rates of Ar and H2 were lowered to 60 sccm, and Ar-4% H2 was used as the reducing gas, rather than H2. As a result, while the in situ XRD result trends are applicable to the bulk foam redox behavior, the precise cycling times for bulk foams are not directly measured by in situ XRD. Attorney Docket No.: 616146.100555 Waterfall XRD spectra of the first two cycles are shown in FIG.4, (a), with diffraction peaks labeled. The corresponding normalized peak intensities, calculated from the peak area of the strongest, non-overlapping peak for each species, for each cycle are shown in FIG.4; each peak is normalized against its own greatest value, rather than against the overall maximum value to better observe the evolution of each phase individually. At the beginning of the first redox cycle, the alloy includes a-Fe (with 1.3%W in solid solution) and µ-Fe7W6. During oxidation, FeWO4forms first, followed by Fe3O4after 10 min. This oxidation cascade from Fe to FeO to Fe3O4 is relatively rapid, with completion taking 20 min after the initial Fe3O4 is formed. The W oxidation rate is however slower (65 min); the discrepancy in oxidation rates leads to the formation of a bcc-W phase between 20 and 70 min, which then slowly oxidizes to FeWO4. During the subsequent reduction, Fe3O4reacts faster, reducing back to Fe in 25 min. By contrast, FeWO4 reduces more slowly, and the final composition includes a-Fe and λ-Fe2W. The reduced Fe2W is nanocrystalline (as discussed previously), and the corresponding diffraction peaks are very broad. As a result, the broadened peak region overlaps with the sharp crystalline peaks for W and FeO; the normalized peak integrals for these two phases were masked during data processing so that erroneous peak identification is avoided. After the first redox cycle, the reaction kinetics are greatly accelerated. This is attributed to two factors: (i) the microstructure, altered by the CVT reduction mechanism, has much higher surface area due to submicron pore formation, and (ii) the Fe2W phase oxidizes more rapidly than Fe7W6, and does not form a large amount of transient W phase during oxidation. In the second cycle, the oxidation rate is slightly faster for FeWO4than for Fe3O4, and both proceed much more rapidly than in the first cycle, with oxidation complete after 15 min (compared to 68 min in the first cycle). During reduction, Fe3O4 reduces first, followed, after 20 min, by reduction of the more stable FeWO4, though the reduction rate is similar for both species. The total reduction time of the second cycle is 25 min, compared to 40 min in the first cycle. After reduction, the Fe2W + Fe phase composition is recovered. Replicate testing across four additional foams showed that the second oxidation was, on average, 55% (±13%) faster than the first oxidation, and the second reduction was, on average, 31% (±19%) faster than the first reduction. All replicates showed the same trend of pure W formation during the first oxidation, and much lower amount of pure W formation in the second oxidation. Microstructural evolution of freeze-cast foams during redox cycling: The microstructure of foams after 1, 10, 20, 50, and 100 cycles is shown in FIG.5 in cross-section micrographs in the reduced (left) and oxidized (right) state. In the reduced state, each lamella includes 2-5 µm ligaments interpenetrated by open porosity (red arrows) of the same size. Each ligament, in turn, includes an Attorney Docket No.: 616146.100555 interpenetrating network of α-Fe (with 1.3%W in solid solution) and submicron λ-Fe2W grains and submicron CVT pores. This same reduced microstructure is maintained through 50 cycles with the microscale porosity increasing as cycling progresses. Since each reduction and oxidation half-cycle fully reduced or oxidized the Fe and W, these cycles represent 100% Fe utilization. The oxidized state includes two interpenetrating networks of Fe3O4 and FeWO4. Between 1 and 10 cycles (FIG.5, (b) and (d), this double network fills most of the microporosity present in the lamellae, but as cycling progresses to 20 cycles (FIG.5, (f)) and beyond, the volume fraction of microporosity increases and the oxide expansion no longer fills all the porosity available in the lamellae. Simultaneously, the FeWO4 domain size (green arrows), which is initially large (up to 10 µm), decreases with cycling as the CVT mechanism redistributes W within the lamellae. After 20, 50, and 100 cycles, the oxidized microstructure closely resembles that of the reduced microstructure, albeit without the submicron CVT pores which are always filled by volume expansion during oxidation, then re-created during the subsequent reduction. The increase in microporosity observed after more extensive cycling is accompanied by a drop in the domain size of each oxide. After the first cycle, the Fe3O4 and FeWO4 domains are ~10 µm in size, with minimal interpenetration between neighboring areas. After 10 cycles, mixing has improved, and the still-distinct domains are ~2-3 µm in size. After 20 cycles (FIG.5, (f)), and further after 50 and 100 cycles (FIG.5, (h)-(i)), the domains are ~1 µm in size, and the FeWO4 domains are evenly distributed within a Fe3O4 matrix rather than segregated in larger regions. Phase composition is confirmed by XRD analysis and discussed below. Remarkably, the oxidized state micrographs indicate that, with increasing cycling numbers, individual lamellae become more porous rather than denser, as would be expected for a high- temperature process due to sintering, and as was observed in previous studies of Fe-Ni, Fe-Cu, and Fe- Co redox systems. Also, the phases become more homogeneously mixed rather than more segregated, as also observed in these other systems. The initially higher density of the lamellae is attributed to the high sintering temperature of 1200 °C. During cycling, the lower cycling temperature of 800 °C is insufficient to induce further sintering in the W-rich phases (Fe2W and FeWO4). Additionally, the high-volume fraction of these phases, combined with the cyclic regeneration of submicron pores by CVT reduction, is sufficient to open new submicron pores that did not initially exist in the as-sintered foam, and to homogenize the size and shape of both new and existing pores. The CVT reduction also serves to break up the initially unmixed FeWO4regions, and redistribute their W content throughout the lamella, leading to an increasingly dispersed microstructure which is expected to be more reactive, due to shorter diffusion distances and increased interfacial area. In addition to new submicron pores Attorney Docket No.: 616146.100555 forming during cycling, the microscale sintering inhibition pores grow in size due to the escape of steam during reduction. The formation of pores during hydrogen reduction is well known from the production of sponge iron, and in the Fe-25W foams these pores do not sinter closed, leading to an increase in pore volume fraction with continued cycling. The major degradation mechanisms (Fe-shell formation and aggregation of the alloying element) observed in previously studied Fe-X redox cycling materials are absent. No shell is observed to form at the edges of the lamellae or at the edges of the foam overall, and no W coarsening is observed. This is probably because, unlike the Fe-25Mo system where MoO2 is formed each cycle, no binary WO2 is observed. Instead, all W content is oxidized to the ternary FeWO4 phase, resulting in atomic mixing of W and Fe in both the oxidized and reduced states. The conclusions drawn from the cross-sectional microstructure are supported by the microstructure of foam surfaces (unsectioned) as well, with representative micrographs after 1, 10, 20, 50, and 100 cycles shown in FIG.6. Initially, the reduced surface contains two levels of porosity: micro-scale, sintering-inhibition pores and submicron CVT pores (FIG.6, (a)). The corresponding oxidation fills much of this porosity, though some micron and submicron porosity remain present (FIG. 6, (b)). Submicron porosity is particularly prevalent in areas with higher W content (indicated in micrographs by lighter Z-contrast). These areas can also show FeWO4whisker formation (FIG.6, (d)), consistent with chemical vapor transport being active during oxidation. With continued cycling, CVT may be responsible for a small amount of mass loss from the foam, measured as ~1 mg per cycle for a Fe-25 at%W (Fe-52 wt% W) foam with 800 mg starting mass; this material condenses onto the sides of the furnace tube and the insulating block during cycling. Analysis of the insulating block and tube inner surface showed significant Fe content as well, indicating that the mass loss measured may instead be due to mechanical spalling of powders during cycling. In either case, although the lost mass is no longer in the foam, it is still redox active and will continue cycling as long as it remains at elevated temperature. Redox cycling again induces homogenization of pore size, pore distribution, and phase distribution within the lamellae. After 10 cycles, the reduced state shows more even distribution of Fe2W, and proliferation of submicron CVT pores. The 10 cycles oxidized state shows that these pores are still largely filled during oxidation, and FeWO4 whiskers are prevalent in some areas, as seen in FIG.6, (d) (green arrow). At 20 cycles in the reduced state, more submicron pores are observed, and the lamellar surface continues to undergo homogenization: rather than a flat Fe sheet dotted with Fe2W grains, the pore homogenization transforms the surface into a 3-D interconnected web of micron-scale Fe ligaments which are prevented from coarsening with each other by Fe2W interconnects. At 20 Attorney Docket No.: 616146.100555 cycles in the oxidized state, significant porosity is present after oxidation, and the oxidized microstructure is similar to the reduced microstructure, except for partial filling of the pores due to the volumetric expansion: the micro scale pores remain largely visible, though the submicron pores are filled. At 50 cycles, the reduced and oxidized microstructure are again similar, showing little change from the 20-cycle state. Some whiskers are still observed in the oxidized state, though they are less prevalent than in earlier cycles. A net decrease in the size of the FeWO4 regions is observed as well, from ~5 μm at 1 cycle, to ~2 μm at 10 cycles, to <1 μm at 20 cycles and beyond (green arrows). Finally, at 100 cycles, the microscale porosity has expanded, while the submicron pores remain intact, with no evidence of degradation as compared to the 50-cycle foam. XRD diffraction patterns of the foams in the reduced state at 1, 10, 20, 50, and 100 cycles, shown in FIG.7, (a), reveal that the phase composition in the reduced state is not changing with cycling, and that the Fe2W XRD pattern contains both a broad (103) peak at 2q = 40-45° (indicative of nanocrystalline structure) and a sharp (110) peak at 2q = 38°. The (200), (112), and (201) peaks at 2q = 44.3, 45, and 45.9°, respectively, are difficult to characterize as crystalline or nanocrystalline due to their proximity to each other and to the Fe peak at 2q = 45.2°: they fall in the shaded region in FIG.7, (a). This coexistence of both crystalline and nanocrystalline peaks in the same diffraction pattern is observed at all cycling points and was also previously observed in our study of Fe-25Mo and by Morales et al. in a study of MoO3 reduction by hydrogen. The similarity of the pattern observed here to the Fe2Mo pattern observed by Morales et al. indicates that the (112) peak observed here is likely crystalline rather than nanocrystalline as well. The oxidized state diffraction patterns (FIG.7, (b)) show crystalline character for both Fe3O4and FeWO4, and the diffraction patterns do not change during cycling, indicative of stable phase evolution and consistent oxidation of both phases. Architecture evolution of freeze cast foams: While the microstructure shows remarkable steadiness, the overall architecture deforms significantly during cycling due to the buckling of the Fe- 25W lamellae. Due to their high aspect ratio, lamellae are easily deformed in the metallic, reduced state, as the expansion and contraction of cycling induce significant stress within the lamellae. These stresses develop because each lamella is constrained by neighboring colonies. Additionally, the high operating temperature lowers the yield stress of the metal, resulting in plastic deformation at lower stresses, as well as creep deformation which can result in additional deformation over long cycling times. The reduced, metallic state is expected to be relatively prone to plastic buckling due to the high porosity of the lamellae. Because the W-containing compounds effectively prevent sintering entirely in the foam, the buckling damage accumulates without sintering of neighboring lamellae. As a result, as shown in FIG. Attorney Docket No.: 616146.100555 8, the foams exhibit severe lamellar buckling in both axial and radial directions. This deformation frequently leads to the formation of buckled bundles, characteristic of buckling of thin objects. The axial buckling is least severe at the top of the foam and grows in amplitude towards the bottom of the foam, resulting in a characteristic brooming deformation (FIG.8, blue dashed lines). This shape may be explained by either gravity or the expansion of the foam against the underlying alumina plate causing buckling along the axial direction. The lower side of the foam, while still porous, loses the directional porosity initially present from freeze casting. The severe plastic buckling does not reduce the redox reactivity of the foam, though it can result in increased tortuosity and loss of macroporosity, particularly at the center of the foam where many buckled lamellae meet, resulting in a loss of the original directional porosity (FIG.8, red circle). Micro- and macrostructure evolution of powder beds: The excellent resistance against degradation inherent to the Fe-25W microstructure and the continued performance of the freeze-cast foams (even after undergoing severe buckling and shape change) both question whether similar performance can be achieved from a packed powder bed without a lamellar architecture. Surface and cross section micrographs of reduced powder beds after 1, 10, 20, 50, and 100 cycles are shown in FIG. 9. The microstructure is identical to that of the freeze cast foams, with the same micron- and submicron porosity observed throughout the foam. The Fe ligaments decorated with Fe2W nanograins are observed as well, and they exhibit the same structural characteristics as in the freeze-cast foams. Furthermore, the redox cycling induces homogenization in the powder bed (as in the foams), even from a relatively coarse starting state; the initially micron-size domains of the Fe2W phase (FIG.9, (a), green arrows) are fragmented into submicron domains (red arrows) interrupted by submicron pores after 10 cycles, and the structure remains well homogenized afterwards, with high porosity and reactivity due to the micron- and submicron porosity networks present. The macroporosity differs however: the packed beds have a much higher tortuosity, with no structured channels for gas ingress and egress built into the architecture, and with a lower initial porosity of 78% (vs.92% for the freeze-cast foams). Some larger channels exist, however, due to the agglomeration of different regions during packing. With continued redox cycling, the bed porosity increases, resulting in easier gas access to all parts of the bed. Conclusions A very high resistance against degradation is observed in Fe-25W (at%) freeze-cast lamellar foams undergoing steam-hydrogen redox cycling at 800 ºC. This resistance is attributed to the inhibition of sintering provided by W, as well as to the phase evolution of the Fe-25W system. Attorney Docket No.: 616146.100555 Tungsten is always found atomically mixed with Fe, either in the intermetallic λ–Fe2W or in the mixed oxide FeWO4 phases. This effectively prevents W from segregating into a pure W phase that would reduce the degradation resistance of the foam, as observed for Mo in the Fe-25Mo system. The lasting redox reactivity of the foam also depends upon its hierarchical porous architecture: (i) macroscopic freeze-cast channels with low tortuosity, allowing easy gas access into, and out of, the foam, (ii) microscale sintering inhibition pores, fully interconnected, which increase the surface area (and thus reactivity) of each lamella, and (iii) submicron chemical-vapor-transport pores, regenerated at each reduction half cycle, which greatly increase surface area and limit damage accumulation by effectively reestablishing the original microstructure after each cycle, producing nanocrystalline λ– Fe2W with corresponding nanopores. This microstructural stability allows for each lamella to grow progressively more porous (rather than denser) with cycling, with the W-rich phases becoming both smaller in morphology and more homogeneously distributed throughout the lamellae, increasing the sintering inhibition effect and generating more submicron pores to accelerate subsequent reaction. During redox cycling, the lamellar structure gradually degrades into a less directional morphology, as neighboring lamellae buckle and contact each other. Unlike previously studied Fe-X systems however, this does not impact the overall reactivity of the foams, as the microscale and submicron pore networks remain intact. Cycling of powder beds show similar characteristics to directional freeze-cast foams, indicating that, for the Fe-25W system, directional porosity is not necessary for long-term cycling stability. EXAMPLE 2: MICROSTRUCTURAL EVOLUTION OF Fe-25W POWDER BEDS DURING CO2-H2 REDOX CYCLING AT 800 ºC This example discloses fixed powder beds of Fe-25W at% powders showing excellent resistance to degradation during CO2-H2 redox cycling at 800 °C, releasing CO and H2O, respectively, with 100% metal utilization. During cycling up to 165 full redox cycles (~ 500 h), the reaction kinetics improve with cycle number and a stable, hierarchically porous microstructure forms that shows neither sintering nor densification. The sintering inhibition provided by W addition to Fe stems from the formation of Fe-W mixed phases: as an intermetallic compound (Fe2W) in the reduced state and as a ternary oxide (FeWO4) in the oxidized state. In addition to the large channels running between coarse (100 µm) powder agglomerates, each agglomerate in the powder bed shows both microporosity due to sintering inhibition, and submicron porosity due to a chemical vapor transport reduction mechanism that is active when FeWO4 is reduced by H2 at high temperature. This hierarchical porosity enables fast Attorney Docket No.: 616146.100555 reactions due to the short diffusion distance to free surfaces, and non-tortuous gas flow through the powder bed. As a result, under high CO2 flow, the CO2:CO ratio is sufficiently high to oxidize Fe3O4 to Fe2O3, improving the CO conversion capability on a per-mole basis. In EXAMPLE 1, we showed the excellent characteristics of Fe-25 at.%W (Fe-52 wt.%W) powder beds under H2O / H2 redox cycling, a very similar process of high-temperature oxidation and reduction. The Fe-25W powder bed shows hierarchical porosity: (i) a continuous microporous network that does not densify due to the sintering inhibition of W; (ii) submicron pores within the struts of the network, formed by the chemical vapor transport reduction of FeWO4. This hierarchical porosity was observed both in freeze-cast foams and in fixed powder beds. Additionally, W does not segregate into pure metallic W or WO2, due to the atomic scale mixing of W with Fe in both the reduced and oxidized states. The nominal overall redox reaction of this system is shown in EXAMPLE 1. In the reduced state at 800 °C, the Fe-25W alloy includes an equimolar mixture of α-Fe and λ-Fe2W. During oxidation by steam (H2O), the alloy is transformed into a two-phase oxide mixture of FeWO4 and Fe3O4(with FeO as a transient intermediate phase) and H2is liberated: ^^^^ + ^^^^2^^ +20 ^^2^^ ↔ ^^^^^^^^4 +2 20 ^^^^3^^4+ ^^2Because, at intermetallic Fe2W or the mixed oxide FeWO4 phase - segregation of W is limited. The overall reaction path is shown in the ternary phase diagram 800 °C isothermal section. In this exemplary study, we explore the behavior of the Fe-25W system when oxidized by CO2, rather than H2O, with the nominal redox reaction as: ^^^^ + ^^^^ ^^ +20 ^^^^ ↔ ^2 20 22 ^^^^^^^4 +^^^^3^^4+ ^^^^ A high (or 0.532 kg CO per kg of Fe-25W alloy) makes this system attractive, even though the high mass fraction of W (52 wt%) strongly increases the density of the alloy as compared to pure Fe. In this exemplary study, we show, using in situ x-ray diffraction and thermogravimetry, that the rate of reaction (which is effectively the CO2utilization rate) increases with redox cycle number due to the formation and expansion of porosity in the powder. Additionally, we identify, for the first time, a Fe-based alloy where Fe is oxidized to Fe2O3 under CO2, going beyond the current limits for unalloyed Fe which only achieves Fe3O4. The combination of high CO2 utilization, microstructural stability, and redox cycling stability demonstrate that Fe-W powder beds are excellent candidates for future large- scale CO2 utilization. Attorney Docket No.: 616146.100555 Methods Powder bed fabrication: Fe-25W (at%) powder beds were fabricated from oxide precursors. Slurries were fabricated with Fe2O3 (20.4 wt%, Noah technologies, < 3 µm, 99.99%) and WO3 (19.7 wt%, US Research Nanomaterials, 100 nm, 99.5%) powders added to DI water (59.9 wt%) and ball milled on a rotary mill for 24 h to achieve good mixing. Each slurry contained 9.07 g solids. The slurry was dried overnight in a fume hood, then ground into coarse powder agglomerates (approximately 100-300 µm in size) with a mortar and pestle. This coarse powder was poured into an alumina tube (9 mm ID) with quartz wool caps and exposed to flowing H2, first at 600 °C for 4 h to complete reduction to the metallic state, then at 1200 °C for 3.5 h to achieve partial sintering and give the powder bed sufficient strength to be handled while maintaining porosity throughout the bed. The same procedure was used to fabricate the unalloyed Fe powder beds from the Fe2O3powder, but without the sintering step at 1200 ºC, since reduction at 600 °C was sufficient to partially sinter the pure Fe beds. Heating and cooling rates of 10 and 5 °C / min, respectively, were used throughout. After sintering, each Fe-25W powder bed was sectioned with a razor into individual cylindrical bodies with a mass of ~1 g each, with 7 mm diameter and 10 mm length. Redox cycling: Redox cycling was carried out in an alumina tube furnace (20 mm ID, 500 mm length) at 800 °C (10 °C / min heating rate) under flowing H2and CO2. Alternating reduction and oxidation half-cycles, 90 min each, were achieved by delivery of pure H2 (UHP, Airgas) and pure CO2 (Industrial grade, Airgas) using mass flow controller at 200 and 60 sccm, respectively. The gas used during cooldown was either 100 sccm H2(for powders in their reduced state) or 20 sccm Ar-4% H2(for powders in their oxidized state). A cooling rate of 5 °C / min was used, and the powder beds were brought to room temperature before removing from the furnace. During cycling, the powders were unconstrained, and thus free to expand and contract during the redox cycles. Powder bed microstructural characterization: To obtain cross-sections, powder bed bodies were mounted in epoxy and vacuum-infiltrated to fill porosity before grinding and polishing to 1 µm. Polished samples were sputter-coated with 9 nm Au / Pd for electron microscopy characterization, which was performed with Hitachi SU8030 or Hitachi S4800 instruments on either unmounted beds (with as- fabricated surfaces) or mounted and polished bed cross-sections. X-ray characterization: In-situ X-ray diffraction spectra recorded during redox cycling of powder specimens were collected in a similar setup to our previous work, with minor modifications. Spectra were collected on a Stadi-MP (Stoe, Germany) instrument, with an asymmetric curved Ge monochromator under pure Ag-Ka1 radiation (λ = 0.56 Å) and a one-dimensional silicon strip-detector (MYTHEN2 1k, from Dectris, Switzerland), operated at 40 kV and 40 mA (Beam Size 4 × 0.8 mm). Attorney Docket No.: 616146.100555 Data were collected in Debye-Scherrer (transmission) geometry, using 1-min scans with 2θ diffraction angles spanning 5.7 to 24.3º, after calibration against a NIST Si standard (640d). A portion of the powder bed was placed in a 1.5 mm diameter quartz capillary and surrounded on both ends by quartz wool. The remainder of the capillary volume on the outlet end of the sample was filled with porous ceramic blocks to prevent sample movement upon changes in gas flow. The capillary was installed into a water-cooled, graphite-heated furnace, with temperature stability of 0.1 °C. The experiment began by heating the sample to 800°C (30°C / min) under 80 sccm flowing Ar- 4% H2, to ensure the sample remained metallic until the cycling temperature was reached. Once at 800°C, the gas stream was switched to CO2 for oxidation, also at 80 sccm. Following complete oxidation, noted by unchanging diffraction patterns, Ar-4% H2was flowed at 80 sccm for reduction. This diluted reduction gas was used to slow the reduction reaction to match the time-resolution of the data collection. A second experiment tested a powder bed sample that had been cycled in the redox station for 10 cycles, then brought to the in situ XRD station and tested under the same conditions as above, but at 900 °C. Processing of diffraction patterns was done in Python, with waterfall plots of diffraction spectra, after background correction with a modified polynomial fit. For each phase of interest, the strongest non-overlapping diffraction peak, indexed using reference patterns from the Inorganic Crystal Structure Database (ICSD), was fitted to a Pseudo-Gaussian model using the lmfit package. These fitted peaks were then integrated to achieve an area, assumed to be proportional to the volume of the diffracting phase, and then normalized to the respective maxima of each phase. Thermogravimetry: Thermogravimetric thermal analyses were performed in a Netzsch STA 449 F3 Jupiter Simultaneous Thermal Analysis (STA) instrument.20 mg of each powder sample was pre- cycled to the desired cycle number (0, 10, 100, or 165) and was placed in an alumina crucible with weight of 200 mg. Each sample was measured under 30 sccm CO2 (Research grad, Airgas) with a cover gas of 25 sccm helium (UHP, Airgas). Each sample was ramped and held at 800 °C for 4 h before cooling to room temperature, with heating and cooling rates of 20 and 40 °C / min, respectively. Buoyancy effect for the gas conditions was corrected by measuring the empty crucible under the same measurement conditions used for the samples Performance of the thermobalance of the STA was verified by using a certified sample of calcium oxalate monohydrate (European Pharmacopoeia Reference Standard) up to 1000 °C. To normalize the data run-to-run, a calibration run with no sample was conducted, and all masses were normalized to the recorded stabilization of gas-flow mass gain (at t = 30 min) from the control run. Results and Discussion Attorney Docket No.: 616146.100555 Phase evolution of Fe-25W powder beds: FIG.10, (a) shows time-resolved diffraction patterns for the first two CO2-oxidation / H2-reduction cycles at 800 °C. FIG.10, (b)-(c) shows calculated normalized peak integrals for the strongest peak of each phase, with metallic species shown in FIG.10, (b) and oxide species shown in FIG.10, (c). No carbide phases are observed at any point. Initially, the powder bed includes a two-phase mixture of µ-Fe7W6 and α-Fe(W), which are the equilibrium phases at the initial sintering temperature of 1200 °C, with 1.3 at% W in solid solution within α-Fe(W). When first exposed to CO2at 800 °C (at t = 3 min), the powders start to oxidize immediately. Based on the ternary Fe-W-O phase diagram isothermal section at 800 °C, oxidation is expected to occur in three distinct steps: (i) oxidation of l-Fe2W to FeWO4 + Fe (ii) oxidation of residual Fe to FeO and (iii) oxidation of FeO to Fe3O4. However, since the powder bed initially contained metastable µ-Fe7W6formed at 1200 °C rather than l-Fe2W stable at 800 ºC, the oxidation pattern of the first oxidation differs from the phase diagram. The FeWO4 phase is first observed at t = 4 min and takes ~60 min to fully form. The formation of FeWO4occurs in two stages, as shown from the change in slope of the normalized peak integral in FIG.10, (c) (green). Before t = 18 min, FeWO4 forms relatively quickly, and no other oxide phases are observed. For the next 12 min (t = 18-30 min), both FeWO4 and FeO continue to form at approximately the same rate. At t = 30 min, several changes occur: FeO begins to be further oxidized to Fe3O4, an α- W peak is observed, and the rate of formation of FeWO4 slows significantly. FeO oxidizes to Fe3O4 over a 45 min period (t = 18-63 min), and the relatively sluggish oxidation of α-W into FeWO4 occurs over a 30 min period (t = 50-80 min). Overall, the first CO2oxidation half-cycle takes 75 min to complete, and the final state includes a mixture of FeWO4and Fe3O4, as expected. The first H2 reduction half-cycle is comparatively quick, starting at t = 83 min. First, Fe3O4 rapidly reduces to FeO, over a 2-min period. Then, FeO reduces to Fe over a 10-min period. Reduction of FeWO4to λ-Fe2W, the equilibrium intermetallic phase at 800 °C, is the final and slowest step, occurring over a 20-min period. This reduction step is expected to occur by the chemical vapor transport reduction mechanism discussed above: FeWO4 is converted to a transient vapor phase WO2(OH)2, which condenses under flowing H2on a freshly reduced metallic Fe surfaces nearby, where it forms the intermetallic λ-Fe2W. Because of this reduction mechanism, the λ-Fe2W formed during cycling is nanocrystalline, as evidenced by broad, low-intensity diffraction peaks. Because of the broad shape of these peaks, the Fe2W normalized integral signal (black) is low compared to the initial µ- Fe7W6peak or the α-Fe peak, though the amount of W in the system overall is the same. The second CO2 oxidation half-cycle, beginning at t = 108 min, is significantly faster than the first due to both the different phase composition (i.e., absence of µ-Fe7W6) and the nanocrystalline Attorney Docket No.: 616146.100555 nature of the λ-Fe2W present. The same sequence of oxidation is observed, but it occurs over just a 20 min period, rather than the much slower 70 min period of the first oxidation. No transient W peak is observed in the second oxidation, and the normalized peak integral for FeWO4 (FIG.10, (b)) indicates that the FeWO4fully and rapidly forms before FeO is observed, as expected from the Fe-W-O ternary phase diagram. Finally, the second H2 reduction period is very similar to the first, again lasting 20-min and following the same sequence, including the formation of nanocrystalline λ-Fe2W. Comparison with pure Fe: The Fe-25W powder microstructure is compared with the Fe powder microstructure before (FIG.11, (a)) and after (FIG.11, (b)-(c)) oxidation with CO2 at 800 °C. The unalloyed Fe powder bed (FIG.11a) is initially highly porous, due to the small size of the original Fe2O3powders and the relatively low reduction temperature. Two regions of the oxidized surface are shown in FIG.11, (b)-(c). FIG.11, (b) shows a region of incomplete oxidation, where the growing oxide layer (white dashed region) is growing inwards towards a metallic Fe region (pink dashed region). The oxide grains are much larger than the initial Fe powders, due to both expansion (110% molar volume expansion) and sintering at high temperatures. In other regions (FIG.11, (c)) the oxide surface is completely densified, trapping unreacted Fe below the oxide layer and resulting in incomplete reaction. These thoroughly sintered oxide regions contrast starkly with the porous structures observed, even in the oxidized state, for Fe-25W powders (FIG.11, (e)-(f)). The evolution of Fe-25W powder beds during oxidation was further explored with TGA under flowing CO2, shown in FIG.12, for powder beds either after reduction and sintering (0 cycle), or pre- cycled to 10, 100, and 165 cycles; for comparison, an Fe-only powder was also subjected to TGA after reduction and sintering (0 cycle). A constant flow rate of 30 sccm CO2was applied to each sample during a 20 °C / min ramp to 800 °C, followed by a 4 h hold at 800 °C. The expected mass gain for complete conversion of 3Fe to Fe3O4 (138%, brown dashed line) and for complete conversion of Fe- 25W to Fe3O4+ FeWO4(130%, grey dashed line) are indicated. The TGA data show that the Fe-25W powder bed is sluggish to react during the first cycle, reaching only 24% mass gain in 4 h, compared to an expected maximum mass gain of 30%. The Fe- 25W powder bed is in fact slower to react than the Fe-only powder bed up to t = 180 min, which reaches 23% mass gain in the same period, out of an expected maximum mass gain of 38% for full conversion to Fe3O4. The microstructure of the unalloyed Fe powder after CO2 oxidation in the TGA, shown in FIG.11, shows the formation of a dense oxide shell, and an initially present microporous network that is completely filled with oxide, slowing further reaction, as expected. After 10 cycles, however, the Fe-25W powder bed is much more reactive, rapidly reaching 28% mass gain in the first 40 min, between t = 40 min (T = 575 °C) and t = 80 min (T = 800 °C). This is Attorney Docket No.: 616146.100555 followed by a very slow mass increase to 31% by t = 240 min. The increase above the expected mass gain is likely due to the formation of Fe2O3 in the bed, as confirmed by ex-situ XRD. The mass gain curves for the 100 and 165 cycle specimens are very similar to the 10thcycle one, characterized by a rapid mass gain for t = 40-80 min, followed by a much more gradual mass gain for the remainder of the hold. All three multi-cycle Fe-25W bed samples oxidize beyond Fe3O4 to Fe2O3, confirmed by ex situ powder XRD of the TGA samples. Finally, to assess the formation of Fe2O3under CO2oxidation conditions, in situ XRD redox cycling was performed on a bed that had been pre-cycled to 10 cycles. For this experiment, the cycling temperature was raised to 900 °C to increase the rate of reaction. The results are shown in FIG.13, with stacked diffraction patterns in FIG.13, (a) and normalized peak integrals in FIG.13, (b)-(c). Pure CO2and H2were used as the oxidizing and reducing gases, respectively. The oxidation for the 11thcycle proceeds by the same path and reaction times as seen in cycle 2 in FIG.10. The initial mixture of λ- Fe2W and α-Fe(W) first oxidizes to FeWO4, fully forming after 5 min. Transient FeO is formed at the same time, which undergoes a slower transition to Fe3O4over 7 min. At t = 10 min into the oxidation, as observed in FIG.13, oxidation would be considered complete and H2 would be introduced to begin reduction. If CO2 conditions are held, however, a gradual oxidation of Fe3O4 to Fe2O3 is observed, occurring over a 40 min period (t = 20-60 min). The CO2oxidative conditions are maintained for another 80 min, but no further changes to the phase composition were observed. The slow rate of oxidation indicates that the formation of Fe2O3 observed is indeed due to CO2, rather than an air leak or other unintended oxygen exposure, which would cause very rapid (< 1 min) oxidation to Fe2O3. The formation of Fe2O3is attributed to the combination of a highly reactive bed microstructure subjected to a high flow rate of pure CO2. Such conditions, though not usually achieved, can produce Fe2O3 under CO2 conditions so long as the CO2: CO ratio is very high (< 20 ppm CO). A repeat experiment with lower flow rate (40 sccm) of CO2did not show Fe2O3formation, confirming that a high partial pressure of CO2 is needed for Fe2O3 formation. Subsequent reduction under H2 occurs rapidly, forming first the expected two-phase composition (λ- Fe2W + α-Fe(W)) after 2 min, which then rapidly transforms into a mixture of Fe3W3C and α-Fe(W). Because Fe2W is observed in only one diffraction pattern, it is omitted from the normalized peak integral calculations. The carbide formation is indicative of the presence of solid C in the powder. This carbon is present due to the Boudouard reaction: 2^^^^(^^) → ^^(^^) + ^^^^2 (^^). Theoccurrence of this reaction is attributed to the combination of the reverse gas shift reaction (^^2(^^) + ^^^^2(^^) → ^^2^^(^^) + ^^^^(^^)) when H2 is introduced into the CO2 environment, combined with surfaces onto which C can readily deposit. A repeat experiment Attorney Docket No.: 616146.100555 with a N2flush between oxidizing and reducing conditions did not show carbide formation, indicating that the reverse water-gas shift reaction is responsible for the carbon deposition and carbide formation observed in FIG.11. The carbide phase can be re-oxidized, as shown in the second oxidation starting at t = 140 min. This releases the CO that was originally disproportionated in the Boudouard reaction. A final reduction at t = 160 min again shows carbide formation. Carbide formation could be suppressed by reacting at higher temperatures, shifting the Boudouard equilibrium to further favor CO, and future work will detail the effect of higher temperature on the reaction process. Microstructural evolution: The microstructural evolution of powder beds undergoing CO2 / H2 cycling is remarkably similar that of H2O / H2cycling, since the phases formed are the same in both cases (provided carbide formation is avoided through high flow rate of CO2). Examples of the porous surface microstructure after 1, 10, 100, and 165 cycles in both the reduced and oxidized states are shown in FIG.14. After 1 cycle, the reduced state (FIG.14, (a)) shows a two-phase composition of α- Fe(W) and λ-Fe2W. The α-Fe(W) regions are 2-5 µm in size (larger regions marked with cyan arrows) and relatively dense, while the λ-Fe2W regions are < 1 µm in size and contain submicron porosity (red arrows). This submicron porosity forms due to the chemical vapor transport reduction mechanism. Throughout the specimen, microscale porosity is prevalent (orange arrows). In the oxidized state (FIG. 14, (b)), however, the volume expansion associated with oxidation results in most pores being closed in the first cycle, leaving a mostly densified oxide surface. After 10 cycles, the microstructure of the powder bed is quite different, with a foam-like architecture, as shown in FIG.14, (c) for the reduced state. The large α-Fe(W) regions have shrunk and homogenized, with no clear segregated regions apparent, indicating both phases are in the micron- or submicron size range. Micropores (orange arrows) are more prevalent, and they are consistently 1-2 µm in diameter. Submicron pores (red arrows) are seen in all ligaments of the foam, as expected from the nanocrystalline λ-Fe2W phase formed during H2 reduction. The 10-cycle oxidized state (FIG.14, (d)) shows that the micropores are now present even after oxidation (orange arrows). The submicron pores have closed, but the combination of micropores and homogeneously-mixed FeWO4and Fe3O4prevents sintering or segregation from occurring. After 100 cycles, the microstructure continues to homogenize and become more foam-like. The ligaments of the foam, shown in the reduced state in FIG.14, €, comprise equiaxed particles with no clear sintering or segregation, and ample microporosity. The small size of the particles makes distinguishing submicron pores more difficult, but the inset (red border) shows nano-scale particles with submicron pores, as expected. The oxidized state (FIG.14, (f)) remains foam-like, showing Attorney Docket No.: 616146.100555 similar particle size, and the inset (red border) shows that submicron pores (red arrow) are present in the oxidized state as well, though each phase’s size is microcrystalline rather than nanocrystalline. There are no significant changes to the microstructure after 165 cycles, in either the reduced state (FIG. 14, (g)) or the oxidized state (FIG.14, (h)). In each, micropores and submicron pores are still observed throughout the powder bed. FIG.15 shows representative cross sections of the macro-, meso-, and microstructure after 0, 10, 100, and 165 cycles, in the reduced state. The macrostructure of the powder bed changes significantly over 165 cycles, as shown in FIG.15, (a), (d), (g) and (j). After the initial sintering step (uncycled), each mm-scale powder bed portion contains a wide distribution of particle agglomerate sizes and porosities, with 20 µm-scale channels (yellow letter C) present between 100 µm-scale powders agglomerates. This structure is also seen after 10 cycles (FIG.15, (d)), with no significant changes. The wider channels (yellow letter C) between agglomerates seen after 10 cycles is due to the powder bed expanding during mounting in epoxy. After 100 and 165 cycles, sintering of neighboring agglomerates results in the bed being much more homogeneous, with no distinct particle agglomerates observed and uniform microporosity throughout. The overall mesostructure of the powder bed also becomes more porous with cycling, as shown in FIG.15, (b), (e), (h) and (k). Initially, each powder agglomerate in the bed (FIG.15, (b)) shows significant (65%) microporosity due to the sintering inhibition effect of W. The microstructure of the cross section (FIG.15, (c)) shows ligaments composed of 2-5 µm size regions of α-Fe(W) (blue arrow) and λ-Fe2W (orange arrow) phases. After 10 cycles (FIG.15, (e)) the mesostructure is similar to the uncycled state, but the microstructure (FIG.15, (f)) shows smaller, 1 µm size regions of Fe (blue arrow) and λ-Fe2W (orange arrow) in the ligaments, with submicron pores present as well. After 100 cycles (FIG.15, (h)), the porosity of the mesostructure has greatly increased to 80%, while the microstructure (FIG.15, (i)) shows submicron phases and a higher volume of submicron pores, with the Fe (blue arrow) and λ-Fe2W (orange arrow) regions smaller and more difficult to distinguish. After 165 cycles, both mesostructured (FIG.15, (k)) and microstructure (FIG.15, (l)) closely resemble those of the 100-cycle bed. The overall increase in porosity with cycling is attributed to the cyclic expansion of existing pores by steam released during each reduction half cycle, similar to the formation of pores in the H2-reduction of iron oxide pellets, coupled with a near-total inhibition of sintering and densification. These two processes together result in an increase in porosity in the bed mesostructure. Conclusions The microstructural evolution of Fe-25W powder beds is studied during high temperature CO2 / Attorney Docket No.: 616146.100555 H2redox cycling. After a sluggish first step, the beds show accelerating redox kinetics, particularly for oxidation, attributed to the increase in porosity and decrease in phase size that occurs due to the chemical vapor transport reduction of the mixed oxide FeWO4, forming submicron Fe2W each cycle. The microstructure of the powder beds becomes increasingly porous with cycling due to this chemical vapor transport reduction in combination with the excellent sintering inhibition of the intermetallic Fe2W and the mixed oxide FeWO4. The increasing porosity and decreasing particle size of the powder bed allow for Fe2O3to form under sufficiently high CO2flow rate at 800 °C, increasing the oxygen capacity of the material. Overall, this study demonstrates that Fe-25W powder beds are promising candidates for CO2 utilization, particularly when used for numerous redox cycles. EXAMPLE 3: TUNGSTEN'S ROLE IN ENHANCING SINTERING RESISTANCE OF Fe-W HIERARCHICAL FOAMS DURING REDOX CYCLING This example discloses directional freeze-cast Fe-W lamellar foams with 10‒33 at% W showing distinct microstructural evolutions during steam / hydrogen redox cycling between oxidized and reduced states at 800 ⁰C, depending on W concentration. The Fe-18W and Fe-25W foams exhibit a sufficient volume fraction of W-rich phases – λ-Fe2W to inhibit sintering for α-Fe in the reduced state and FeWO4 to inhibit sintering for Fe3O4 in the oxidized state – thus forming ligaments comprising two phases (Fe / λ-Fe2W and Fe3O4 / FeWO4, respectively). In contrast, a Fe-10W foam with a lower volume fraction of W-containing phases (λ-Fe2W and FeWO4) shows lamellae densification as well as core- shell structure formation, due to Fe outward diffusion during oxidation. While higher W concentration enhances the stability of lamellar structure in Fe-W foams, degradation still occurs, via buckling of lamellae and swelling of foams after extensive cycling. In situ XRD characterization shows that W addition has a minor effect on the oxidation process but slows reduction due to the sluggish kinetics of FeWO4 reduction. This influence is mitigated by the formation of nanocrystalline W-rich phases due to the chemical vapor transport (CVT) mechanism during the reduction of FeWO4 to boost the reaction kinetics during redox cycling. Specifically, we investigate the role of W concentration on the structural evolution of Fe-W foams during high-temperature redox cycling, by varying the W concentration from 10 to 33 at% compared to a single composition of Fe-25W in the previous study. This introduces different volume fractions of W-rich phases (λ-Fe2W and FeWO4) acting as sintering inhibitors in the reduced and oxidized states during redox cycling, thus revealing the critical roles of W concentration on foam degradation. In situ XRD measurements are performed on Fe-10W, Fe-18W, and Fe-33W to Attorney Docket No.: 616146.100555 understand the effect of W concentration on redox kinetics. The structure of low- and high-W foams after various cycles are characterized from micro- to macroscale to reveal the distinct evolution processes during redox cycling with steam and hydrogen at 800 ⁰C. The mechanisms of different evolution processes of lamellae structures in low- and high-W foams are compared to assess the impact of W concentration on the sintering resistance of foams. A compositional window for designing Fe-W foams with high degradation resistance is suggested, enabling Fe-based foams for long-term redox cycling for energy storage. Methods Freeze casting, reduction, and sintering of Fe-W foams: Fe-W foams with four compositions, i.e., Fe-10W, Fe-18W, Fe-25W, and Fe-33W (at%), were prepared by freeze-casting of aqueous suspensions, ice sublimation, hydrogen reduction, and sintering, via methods previously reported for Fe foams and Fe-based foams (Fe-Ni, Fe-Co, Fe-Cu, Fe-Mo). In brief, aqueous suspensions were prepared by mixing Fe2O3powders (Noah Technologies, 99.9%, < 3µm), WO3nanopowders (SSNano, 99.5%, < 100 nm), DI water, and Zephrym PD 4974 (Croda) as a dispersant. Volume fractions of precursors in suspensions for various Fe-W foams are listed in Table 2. The suspensions were milled with yttria- stabilized zirconia milling balls for 24 h on a roller mixer. After milling, 2.0 vol% polyethylene glycol (PEG, Mn=3350, Sigma Aldrich) was added as a binder. The suspensions were vortexed for 2 min and then degassed and cooled in an ice-water bath (~0 ºC). Directional freeze-casting was performed on a copper plate cooled by a thermoelectric chiller (Mauser Electronics). The suspensions were poured into a hollow Teflon mold (outer diameter: 25 mm, inner diameter: 15 mm, height: 15 mm) placed on the copper plate. The plate was then cooled to -30 ºC following an exponential cooling function, so as to initiate ice growth at a constant rate from the bottom to the top of the mold. This results in ice dendrites with constant thickness and spacing throughout the height of freeze-casted foams. After demolding, a 1-2 mm high section at the bottom of the frozen specimen was cut and discarded. The cast specimens were placed in a freeze drier at a temperature of ˗54 ºC under vacuum (~13 Pa) for 24 h to sublimate the ice completely. After ice removal, lamellar green bodies were thermally treated to achieve metallic foams in three steps under ultra-high purity (UHP) H2 (Airgas) flow: (i) debinding at 300 ºC for 1 h to burn out all organic components; (ii) reduction of Fe2O3and WO3at 600 ºC for 4 h; (iii) sintering at higher temperatures for various time, depending on W concentrations, to partially densify lamellae. For the four Fe-W foams, thermal profiles of the debinding and reduction steps and cooling rates (10 ºC / min) were the same, while temperatures and times for sintering (listed in Table 2) were varied to achieve Attorney Docket No.: 616146.100555 similar porosities in foams with different compositions. The morphologies of as-sintered Fe-W foams with various W concentrations are shown in FIG.34. Table 2. Precursor volume fraction in suspensions and peak sintering profiles for the four Fe-W foams studied. Foam Fe2O3WO3DI Water Zephrym Polyethylene Peak sintering (at%) (vol%) (vol%) (vol%) (vol%) glycol (vol%) Fe-10W 8.08 1.92 87.49 0.51 2.00 1000 ºC / 3.5 h Fe-18W 6.81 3.18 87.48 0.53 “ 1100 ºC / 3.5 h Fe-25W 5.84 4.15 87.46 0.55 “ 1200 ºC / 3.5 h Fe-33W 4.87 5.12 87.45 0.57 “ 1000 ºC / 7.0 h Redox cycling of foams: Redox cycling of metallic foams was conducted in an alumina tube furnace (outer diameter: 25 mm, inner diameter: 20 mm, length: 500 mm) at 800 ºC under flowing steam or UHP H2. During the oxidation process, steam was supplied by flowing UHP Ar (Airgas) at 120 sccm, controlled by a mass flow controller (MKS Instruments), through a water bubbler heated to 93 ºC to achieve an H2O partial pressure of 77 kPa, enabling oxidization of Fe into Fe3O4rather than Fe2O3. During the reduction process, UHP H2 was supplied at 200 sccm directly by another mass flow controller. A single redox cycle includes both oxidation and reduction processes, a duration of 90 min for each, ensuring complete redox reactions (100% Fe utilization). After redox cycling, specimens were cooled to ambient temperature within the tube furnace by flowing either 20 sccm Ar-4% H2gas for oxidized foams or 100 sccm H2 for reduced foams. Microstructure characterization: Foams were cold mounted with epoxy resin (Epothin 2 Resin, Buehler), vacuum infiltrated to fill porous structures, cured for 24 h, and polished with diamond suspensions (MetaDi, Buehler) down to 0.25 μm. For metallic foams in their reduced state, oil-based suspensions were used to avoid corrosion during polishing. Specimens were coated with a 6 nm thick osmium conductive layer and examined by scanning electron microscopy (SEM) using a Quanta 650 ESEM (ThermoFisher Scientific) instrument equipped with an Oxford Aztec Energy dispersive X-ray spectroscopy (EDS). To evaluate surface areas in as-sintered Fe-W foams with various of W concentrations, Fe-10W, Fe-18W, and Fe-25W foams were measured by a Brunauer-Emmett-Teller (BET) method. Nitrogen adsorption of foams was performed using a Micromeritics 3-flex apparatus at -196°C. Prior to analysis, samples were degassed stepwise up to 300°C for 8 h under vacuum. The degassing conditions have Attorney Docket No.: 616146.100555 three steps: (1) Heating the foam sample at 100°C for 60 minutes under vacuum to remove moisture and volatile impurities; (2) Increasing the temperature to 200°C and hold for 120 minutes to eliminate any remaining moisture and adsorbed gases; (3) Raising the temperature to 300°C and maintain for 240 minutes to remove strongly adsorbed impurities. Nitrogen adsorption isotherm covered a range of relative pressures (p / p°) from 0.05 to 0.30. X-ray diffraction characterization: In situ X-ray diffraction (XRD) was performed on freeze-cast Fe-W foams to characterize phase evolution during redox cycling in a modified Stadi-MP (Stoe, Germany) instrument equipped with an asymmetric curved Ge monochromator under pure Ag-K^^1 radiation (^^ = 0.56 Å) and a 1D silicon strip-detector (MYTHEN21k, from Dectris, Switzerland), operated at 40 kV and 40 mA (beam size: 4 × 0.8 mm). XRD patterns were collected in Debye– Scherrer transmission geometry, using 1 min scans with 2θ diffraction angles spanning 5.7–24.3°, after calibration against a NIST Si standard (640d). A colony of lamellae was cut from Fe-10W, Fe-18W, and Fe-33W foams in the as-sintered state, with a size of ~1–1.5 mm in height and ~1–1.5 mm in thickness. The specimen was then transferred into a quartz capillary (with an inner diameter of 1.5 mm) and surrounded on both ends by amorphous quartz wool to fix their position. The remainder of the capillary volume, on the inlet and outlet end of the sample, wasfilled with porous ceramic blocks to prevent sample movement upon changes in gasflow. The capillary was inserted into a water-cooled, graphite-heated furnace, with a temperature stability of 0.1 °C. The samples were first heated to 800 °C (heating rate of 30 °C / min) underflowing Ar-4% H2to prevent oxidation prior to redox cycling. After reaching 800 °C, the gas was switched to -UHP Ar, and bubbled through a 40 °C water bubbler for oxidation (PH2O = 7.3 kPa). The gas lines between the bubbler and capillary were heated to >100 °C to prevent steam condensation. After complete oxidation, noted by unchanging diffraction patterns, the gas flow was switched to Ar-4% H2(flow rate of 60 sccm) for reduction. This diluted gas slowed the reduction reaction to match the time resolution of the collection of XRD patterns. The measured XRD data were analyzed in MATLAB to plot the evolution of XRD patterns as a function of time after background correction with a modified polynomialfit. To study the evolution of each phase as a function of time, the strongest diffraction peaks (i.e., (110) reflection for α-Fe, (103) reflection for λ-Fe2W, (1 -20) reflection for µ-Fe7W6, (113) reflection for Fe3O4, and (111) reflection for FeWO4, indexed using reference patterns from the Inorganic Crystal Structure Database (ICSD)) werefitted with a Lorentzian function. Thesefitted peaks were then integrated to calculate their area and then normalized by the respective maximum value. Attorney Docket No.: 616146.100555 Ex situ XRD patterns of samples were measured in a reflection mode with pure Cu Kα1 radiation source (^^ = 1.54 Å) with a Rigaku Smartlab 3kW Gen2. The reflection 2θ angles were scanned in a range of 25–60º at a scanning speed of 2º / min. Results and discussion Phase evolution and thermodynamics during redox cycling - Initial phases: The tungsten concentration in foams was varied in a range of 10‒33 at%, i.e., Fe-10W, Fe-18W, Fe-25W, and Fe- 33W (vertical dashed lines in FIG.16, (a)), to increase mole fractions of the intermetallic phase λ-Fe2W from 12.5 to 100%. The Fe-W foams were sintered at various temperatures between 1000 and 1200 ⁰C for different times, depending on W concentrations (Table 2). This was done to achieve similar mechanical strength and similar inter- and intra-lamellar porosity in the initial state for each foam composition. From the XRD patterns of as-sintered foams (FIG.16, (b)), the main phases are identified as bcc α-Fe (with 1.3 at% W in solid solution) and λ-Fe2W, except for the Fe-25W foam sintered at 1200 ⁰C, for which the μ-Fe7W6 phase is present, consistent with the phase diagram (FIG.16, (a)). In the Fe-33W foam with a stoichiometry of Fe:W = 2:1, the α-Fe, λ-Fe2W, and μ-Fe7W6are observed simultaneously in the as-sintered state. The presence of α-Fe and μ-Fe7W6phase indicates that the formation of λ-Fe2W phase is incomplete with the current sintering profiles, i.e., 7 h at 1000 ⁰C, in the α-Fe + λ-Fe2W field below the peritectic temperature of 1060 ⁰C. This might be attributed to the sluggish diffusion of Fe and W atoms in Fe-W foams with a highly porous structure. Increasing the sintering temperature and extending the sintering time can promote diffusion and reaction to form a complete λ-Fe2W structure, i.e., Fe7W6 + 5Fe → 6Fe2W. However, increasing sintering time may eliminate porous structures and decrease the kinetics of redox cycling. In our exemplary study of Fe- 25W foams in EXAMPLE 1, μ-Fe7W6 is entirely converted into λ-Fe2W, a thermally stable phase (FIG. 16, (a)), after the first redox cycle at 800 ⁰C. In all Fe-W foams with varying W concentrations, α-Fe and λ-Fe2W remain the primary phases involved in oxidation during redox cycling. Phase evolution during redox cycling: In as-sintered foams with various W concentrations (i.e., Fe-10W, Fe-18W, and Fe-33W), the evolution of XRD patterns as a function of time in the first redox cycle is plotted in FIG.17, (a), FIG.18, (a), and FIG.19, (a), respectively, labelling each phase. The corresponding normalized peak intensities of both metallic and oxide phases were determined as a function of reaction time to demonstrate phase evolution semi-quantitatively. During the oxidation stage, the initial phases (α-Fe and λ-Fe2W in Fe-10W and Fe-18W foams) are oxidized to FeO, Fe3O4, and FeWO4. According to the evolution of normalized intensity, FeWO4forms first, followed by FeO and then Fe3O4 in the first 10 mins (FIG.17, (c) and FIG.18, (c)). The intensity of FeO rises before Attorney Docket No.: 616146.100555 declining as the Fe3O4intensity increases, suggesting an oxidation sequence as Fe → FeO → Fe3O4. This result is consistent with the calculated ternary Fe-W-O phase diagram at 800 ⁰C and in situ XRD findings in a previous study of freeze-cast Fe-25W. The thermally stable Fe3O4 and FeWO4 phases achieve their maximum intensities after oxidation for 20 mins. Although µ-Fe7W6is initially present in the as-sintered Fe-33W foams (FIG.19, (b)), the oxidation behavior is similar to other foams with different W concentrations. During the reduction stage, the reduction of Fe3O4progresses as Fe3O4→ FeO → Fe in the first 20‒25 min. The intensity of FeO diffraction peaks first increases and then drops quickly and disappear, indicating fast reduction to Fe, as shown in FIG.17, (c). The FeO → Fe reduction kinetics by hydrogen at 700 ⁰C are typically slower than the Fe3O4→ FeO reduction step, due to the sluggish outward diffusion of oxygen through a dense Fe shell. The fast reduction of FeO in the present foam can be attributed to the higher reduction temperature of 800 ⁰C and highly porous structures with a large surface area. As a comparison, the necessary time for complete reduction of FeWO4 spans from 30 min in Fe-10W foam (FIG.17, (c)) to 75 min in Fe-33W foam (FIG.19, (c)). This indicates faster reduction kinetics of Fe3O4 than the more stable FeWO4, which needs much longer time for complete reduction, especially in foams with higher W concentration (due to a higher volume fraction of FeWO4 formed). In Fe-10W foam (FIG.17, (b)), the two reduction products, α-Fe and λ-Fe2W, form and grow with similar kinetics. During the reduction of Fe-18W (FIG.18, (b)) and Fe-33W (FIG.19, (b)) foams, the µ-Fe7W6 is observed as a transient, non-equilibrium phase before the equilibrium λ-Fe2W phase forms. The λ-Fe2W intensity increases gradually with the decline of µ-Fe7W6intensity, suggesting the reaction of Fe7W6+ 5Fe → 6Fe2W. In Fe-33W foam (Fe:W = 2:1), both µ-Fe7W6and α-Fe phases first form on reduction, but are then depleted to form a single phase λ-Fe2W after reduction for sufficiently long times (~ 75 min). The complete reduction of Fe-W foam with higher W concentrations is much longer due to the slow reduction kinetics of FeWO4. During reduction, the presence of the transient phase µ- Fe7W6, which is thermally unstable at 800 ⁰C according to the calculated Fe-W phase diagram (FIG.16, (a)), prior to λ-Fe2W indicates a gradual alloying process of Fe with W to form the thermally stable λ- Fe2W. The µ-Fe7W6is not observed in Fe-10W during the reduction process (FIG.17, (b)). Its intensity might be too low to be detected due to the relatively lower W concentration. Thermodynamic consideration of Fe-W foams: During redox cycling at 800 ⁰C, Fe and λ-Fe2W phases are stable in the reduced state and Fe3O4and FeWO4are stable in the oxidized state. The corresponding redox reactions and associated ions and electrons exchange follow the reversible reactions below: 3^^^^ + 4^^2^^ ↔ ^^^^3^^4 + 4^^2 (1) Attorney Docket No.: 616146.100555 ^^ ↔ 4^^ 2− −2 + 4^^ + 8^^ (2)^^^^ ^^ +16^^ ^^ ↔ ^^^^^1 16 2 32 ^^^4 +3 ^^^^3^^4+ 3 ^^2(3) progressing In the context of ROBs corresponds to energy recharge. The reaction Equations (1) and (3) show that the metallic phases in foams are oxidized in the discharging mode by capturing oxygen and then reduced in the charging mode by releasing oxygen in metal-air batteries. The reaction equation (3) also indicates the participation of W in redox cycling by forming W-rich phases, i.e., λ-Fe2W in the reduced state and FeWO4in the oxidized state, which prevents W segregation. From an electrochemical perspective, electrons are exchanged between hydrogen ions and molecules during discharge and charge processes, as illustrated in Equations (2) and (4). For 1 mole of Fe+W atoms, both oxygen and hydrogen capacities increase linearly as W concentration increases from 0 to 33 at%, with oxygen increasing from 1.33 to 1.77 mol and hydrogen increasing from 2.67 to 3.55 mol, as plotted in FIG.29, (a)-(b). However, due to the high mass density of W, the oxygen capacity per alloy unit mass decreases near linearly, from 382 to 289 g / kg (a 24% drop), with a concomitant decrease in hydrogen capacity per alloy unit mass, from 47.8 to 36.2 g / kg (a 24% drop). The electron numbers, which is proportional to the charge capacity, increase from 2.57 to 3.40 × 105C as W concentration increases from 0 to 33 at% in the alloy. The theoretical Nerst potentials (EN) of Fe-W foams show a slight decrease from 0.94 to 0.92 V with increasing W (FIG.29, (c)), indicating a negligible effect of W concentration on the theoretical battery voltage for an ROB using Fe-W foams as the energy storage material. Microstructural evolution during redox cycling - Initial microstructures in as-sintered foams: FIG.20 shows the surface of lamellae in Fe-W foams with various W concentrations after reduction and sintering. To achieve sufficient mechanical strength for handling but maintain the porous microstructure for better redox cycling, Fe-W foams were only sintered partially. Microscale sintering- inhibition pores forming a continuous network are observed at the lamellae surface. Submicron porosity between α-Fe and intermetallic phases is also visible in the cross-sections of lamellae (white arrows in an inserted view in FIG.20, (b)). The micron- and submicron-pores, in combination with the freeze- cast channels (~tens of microns) between lamellae (FIG.29) form hierarchical porosity in the freeze- cast foams, which provide a large amount of surface for redox reactions and easy access for H2 / steam transport, enabling fast reaction kinetics. As-sintered Fe-10W, Fe-18W, and Fe-25W after peak sintering in a temperature range of 1000‒1200 °C for 3.5 h have surface areas of 3.35 m² / g, 5.61 m² / g, Attorney Docket No.: 616146.100555 and 6.82 m² / g respectively according to the Brunauer-Emmett-Teller (BET) measurements. This indicates larger porosity in foams with higher W concentrations because of its sintering inhabitation effect. The preservation of such hierarchical porous structures during high-temperature redox cycling is crucial for prolonging cycling durability and sustaining the charging capacity of rechargeable oxide batteries for a long lifespan. The intermetallic phases ‒ λ-Fe2W and μ-Fe7W6 with sizes of 2‒3 µm, with brighter contrast due to W in BSE micrographs ‒ are evenly mixed with the α-Fe phase in as-sintered foams. With increasing W concentration from 10 to 33 at%, the volume fraction of the intermetallic phase increases, as expected. Moreover, the morphology of the intermetallic phase transitions from isolated particles in Fe-10W foam (FIG.20, (a)) to a 3D network in foams with higher W concentration (FIG.20, (b)-(d)). In Fe-33W foam, α-Fe is also observed to distribute randomly between intermetallic μ-Fe7W6and λ- Fe2W phases, consistent with XRD measurement (FIG.16, (b)), due to insufficient time for diffusion to form a complete λ-Fe2W structure during the sintering process. However, it is difficult to differentiate μ-Fe7W6and λ-Fe2W phases in BSE contrast due to a relatively small difference of W concentrations between the two phases. The SE micrograph of an identical region on the lamellae surface also show no difference in grain morphologies between both the W-rich phases (FIG.30). According to the in situ XRD characterization (FIG.19, (b)), μ-Fe7W6is converted to λ-Fe2W, a thermally-stable phase during redox cycling at 800 ⁰C, after the first cycle. Microstructural evolution: The evolution of the lamellar structures during redox cycling was investigated by inspecting microstructures after reduction and oxidation for different cycle numbers in Fe-10W, Fe-18W, and Fe-25W foams. The Fe-33W foams are not investigated since the high faction of λ-Fe2W requires a long reduction time (~75 mins, FIG.19, (c)). FIG.21 shows radial cross-sections of the three foams after 20 redox cycles. The lamellae in Fe-10W foams display a distinct core-shell structure, i.e., a dense Fe-rich shell (orange arrow) with a porous, W-rich core (FIG.21, (a)-(b)), in contrast to homogenous porous structures exhibited by lamellae in Fe-18W and Fe-25W foams. This core-shell structure is not present in the lamellae of Fe-10W foams in the as-reduced state, comprising a homogeneous mixture of α-Fe with λ-Fe2W phases (FIG.20, (a)), before redox cycling. The result indicates, in Fe-10W during redox cycling, a strong outward diffusion of Fe toward the lamellar surface from the lamellar core, resulting in a highly porous core. In contrast, the α-Fe with λ-Fe2W phases are observed to mix homogeneously and form ligaments with sizes of 3‒5 µm (red dashed circle in FIG. 21, (d)) in Fe-18W and Fe-25W foams after 20 cycles. These ligaments connect with neighboring ligaments to form a 3D-interpenetrating network of α-Fe with λ-Fe2W. The α-Fe and λ-Fe2W phases are much finer in size for foams after 20 cycles compared to as-sintered foams (FIG.20). Moreover, the Attorney Docket No.: 616146.100555 channels between lamellae in Fe-10W after 20 cycles are much narrower compared to those in Fe-18W and Fe-25W foams, implying higher sintering and densification in Fe-W foams with lower W concentration. To understand the evolution process of Fe-10W foams during redox cycling, the lamellar microstructures are shown in FIG.22 after 1 and 20 cycles in the oxidation state. After the first cycle, Fe3O4 (yellow arrow) and FeWO4 (red arrow) phases are mixed throughout the lamellae (FIG.22, (a)- (b)). The α-Fe oxidation (Fe → Fe3O4) induces a notable molar volume expansion (~110%) which fills the microporosity widely distributed within the as-sintered foams (FIG.20, (a)). The FeWO4 phase (with higher brightness in BSE micrographs) is found between the faceted Fe3O4 phase (FIG.20, (b)), inhibiting its contacting and sintering during oxidation. After 20 cycles, oxidized lamellae exhibit dense Fe3O4shells (yellow arrows) and porous cores enriched with isolated FeWO4regions, as shown in FIG. 22, (c). This core-shell structure indicates a gradual outward diffusion of Fe during oxidation which form a dense shell after multiple cycles. The imbalance between the outward diffusion of Fe to the lamellar surfaces and the inward diffusion of oxygen results in Kirkendall pore formation in the lamellar core. This phenomenon was reported in previous studies of pure Fe foams, as well as Fe-Co, Fe-Cu, and Fe-Ni foams alloyed with redox inactive elements – unoxidized by steam under standard redox conditions - which form a metallic core encased within a Fe3O4shells after oxidation. This core- shell structure is observed in both oxidized and reduced lamellae in the present Fe-10W foams, which, at the microscopic level, undergo notable densification due to extensive contact and sintering of neighboring lamellae with pure Fe3O4shells during redox cycling. For instant in lamella #2, two thinner lamellae (#2a and #2b) contacted and then sintered together at the top part but left a fork at the bottom marked by a green arrow in FIG.22, (c). In an inclined view of oxidized lamellae (FIG.22, (d)), cracks are observed directly on the side surface of the Fe3O4 shells. These through-cracks provide ingress of steam and egress of H2during oxidation and oxidation, respectively, in the redox cycles. In Fe-18W foams, a different morphology of oxidized lamellae is demonstrated, with a homogeneous mixture of Fe3O4 and FeWO4 phases crisscrossed by microchannels observed in a cross- section view of lamellae after 1 and 20 cycles, as illustrated in FIG.23. Both oxides, which adhere to each other, form ligaments with neither phase separation nor core-shell structure observed after 20 cycles. After 20 cycles, the size of oxide ligaments separated by microchannels decreases to 2‒5 μm, accompanied by notable refinement of the FeWO4phase to submicron size (red arrows, FIG.23, (b)). This is attributed to a chemical vapor transport (CVT) mechanism during FeWO4reduction which produces nanocrystalline λ-Fe2W, in which mixed oxide FeWO4 decomposes to WOx + FeO4-x by forming vapor-phase WO2(OH)2, which is further reduces under flowing H2 to metallic W, releasing Attorney Docket No.: 616146.100555 H2O. However, the W phase was not detected during the reduction of Fe-W foams by using in situ XRD characterization. A possible reason might be the quick recombination of W with surrounding Fe to form the thermally-stable λ-Fe2W phase at 800 ⁰C. This CVT mechanism has also been observed to induce nanocrystalline λ-Fe2Mo in Fe-Mo foams during redox cycling. The adhesion of Fe3O4to the refined FeWO4 within the oxidized micro-ligaments prevents phase separation and sintering of Fe3O4 in oxidized Fe-18W foams. In the reduced Fe-18W foams (FIG.21, (d)), nanocrystalline λ-Fe2W attached to α-Fe phase also inhibits the sintering and coarsening of the metallic micro-ligaments. Since foams with higher W concentration display excellent resistance to phase separation and degradation of their lamellar structure in both the oxidized and reduced states, Fe-18W and Fe-25W foams were further subjected to 50 redox cycles to investigate their long-term microstructure evolution. FIG.24 shows the gradual evolution of lamellar structures in Fe-18W foams, in the reduced state, after 1, 25, 30, and 50 cycles. The left column of FIG.24 shows BSE micrographs at low magnification displaying lamellar buckling with increasing cycling numbers. The higher magnification micrographs on the right column of FIG.24 show a decrease of ligament size and increasing volume fraction of microporosity within the lamellae. The dimension of the ligaments (highlighted by red dashed circles), including α-Fe and λ-Fe2W bonded together, decreases from 5‒10 μm after 1 cycle (FIG.24, (b)) to 2‒ 4 μm after 50 cycles (FIG.24, (h)). This provides additional free surface area and decreases the diffusion distance, thus accelerating the redox reaction. The results in Fe-18W foams are distinct from the densification of lamellae with a core-shell structure observed in Fe-10W foams (FIG.24, (a)) and in previous Fe-Co, Fe-Cu, and Fe-Ni foams as well. In a previous study of Fe-25W foams, increasing porosity in lamellae is mainly attributed to (i) the presence of sintering inhibitor λ-Fe2W that hinders sintering at 800 ⁰C; and (ii) the regeneration and accumulation of submicron porosity induced by the CVT mechanism during reduction. The present results in Fe-W foams with various compositions indicate that strong sintering inhibition and degradation resistance are dependent on W concentrations as well. A threshold W concentration exists between 10 and 18 at%, which introduces a sufficiently high-volume fraction of λ-Fe2W and FeWO4 to inhibit sintering and phase separation. The 3D interpenetrating network of α-Fe / λ-Fe2W in the reduced state and Fe3O4 / FeWO4in the oxidized state appear to be crucial for lamellae to resist degradation during redox cycling. To confirm phase composition after redox cycling, XRD patterns were collected in reduced Fe- 18W foams after 1, 25, 30, and 50 redox cycles, as shown in FIG.25. The results indicate that the phase composition of the cycled foams is identical to that of the as-sintered foams (FIG.16, (b)), including α- Fe and λ-Fe2W. Notably, the (103) reflection of λ-Fe2W changed from a sharp peak in the as-sintered state (FIG.16, (b)) to a broad hump in the 2θ range of 40⁰ ‒ 44⁰ (grey shaded area) after the first redox Attorney Docket No.: 616146.100555 cycle, implying the presence of nanocrystalline grains. However, the sharp (110) reflection still indicates a crystalline structure for the λ-Fe2W phase. The sharpness of (200) and (112) reflections for λ-Fe2W, i.e., for 2θ angles of 44.1⁰ and 44.9⁰, are difficult to assess due to their proximity to the α- Fe (110) (2θ = 45⁰) reflection. The presence of both sharp peak and broad hump for the λ-Fe2W phase might be attributed to the refinement of grain sizes or the special grain morphologies after reduced cycling. According to the microstructure study, redox cycling refines the grain size of the λ-Fe2W phase and may lead to the co-existence of crystalline and nanocrystalline grains. This phenomenon was also reported in the hydrogen reduction of Fe2MoO4 to Fe2Mo and redox cycling of Fe-25W foams. Macrostructural foam evolution: The cross-section of Fe-W foams was examined at a larger scale to assess the structural evolution of the entire foam. FIG.26 shows a radial cross-section view at the edge of Fe-10W foams in the reduced and oxidized states after 20 cycles, displaying the formation of a macroscale shell structure encasing the entire foam. In the reduced state (FIG.26, (a)), sintering and densification during redox cycling are observed in Fe-10W by the decrease in channel width and more contacts between neighboring lamellae as compared to the as-sintered foams (FIG.31, (a)). Additionally, a dense Fe shell (white dashed line), with a thickness of about 15‒20 µm, forms on the surface of the entire foam. The dense Fe macro-shell decreases gas accesses (e.g., H2 ingress and steam egress during reduction) thus affecting reaction speeds. In the oxidized state, a thicker Fe3O4shell is observed on the surface of foams, as shown in FIG.26, (c)-(d). This thick Fe3O4 shell also hinders steam ingress and H2 egress during the oxidation process. Thus, the core-shell structure of Fe-10W foams at the macro-scale encapsulates the foam interior and then blocks gas channels during redox cycling, leading to sluggish reactions and decreased foam utilization. By contrast, Fe-18W foams demonstrate a quite different evolution of their macrostructure after 20 cycles, as shown in FIG.27. The axial cross-section of the entire Fe-18W foam in the reduced state displays extensive lamellae buckling, similar to that observed in a previous study of Fe-25W foams, but without the formation of a shell encapsulating the foam, as observed in Fe-10W. The high aspect ratio of lamellae makes them susceptible to deformation due to internal stresses caused by volumetric expansion and contraction during redox cycling. However, this deformation is constrained due to their proximity to neighboring lamellae, leading to buckling. Additionally, the strength of the lamellae is low at 800 ºC, and creep deformation is likely, particularly after numerous cycles. Furthermore, increasing porosity in lamellae with cycling numbers reduces the yield strength of lamellae and induces higher degrees of plastic buckling. Lamellae buckle isotopically in both radial and axial directions, resulting in a swelling of the foams. Notably, the buckled lamellae at the bottom of the foam (yellow dotted frame) are extensively deflected in the radial direction (blue arrows), characterized as a brooming deformation Attorney Docket No.: 616146.100555 at the bottom of the foam, most probably due to gravity. In the center of foam marked by a white dashed circle, buckled lamellae contact with their neighbors extensively, resulting in sintering and gradual loss of lamellar structure after long-term cycling. Degradation mechanisms of Fe-W foams: Although adding tungsten, a strong sintering inhibitor, to Fe foams improves the stability of the lamellar structure by reducing lamellae sintering and densification, the core-shell structure at both micro- and macro-scale levels is still observed in Fe-10W foams after 20 cycles. In contrast, Fe-18W and Fe-25W foams are more stable, forming a 3D network of binary phase ligaments (α-Fe / λ-Fe2W in the reduced state and Fe3O4 / FeWO4 in the oxidized state), even after 50 cycles. This is consistent with the observation in our preliminary studies in Fe-25W and Fe-20Ni-20W foams made by the same freeze-casting process. The result demonstrates the crucial role of W concentration and the associated effect in preventing structural degradation after extensive cycles number. The detailed mechanism is illustrated through a series of cross-sectional sketches of lamellae in foams with two W concentrations during redox cycling, as shown in FIG.28. In the low-W foams, α- Fe and λ-Fe2W phases are mixed homogeneously in the as-sintered lamellae initial microstructure (FIG. 28, (b)). The phase separation between Fe3O4 and FeWO4, leading to a core-shell structure, is illustrated in the oxidized lamellae after several redox cycles (FIG.28, (c)). As Fe diffuses outwards to the lamellar surface, a dense Fe3O4shell forms, encasing each lamella and decreasing their surface during oxidation at 800 ⁰C. After reduction, the shell structure in lamellae is preserved (FIG.28, (d)), inducing sintering and densification of neighboring lamellae. In the high-W foams, a larger volume fraction of λ-Fe2W is mixed within the α-Fe phase to form a 3D interpenetrating network, as depicted in FIG.28, (e). The adhesion between both phases suppresses the outward diffusion of Fe during oxidation and prevents the formation of detrimental core- shell structure in the oxidized state, as sketched in FIG.28, (f). Additionally, the FeWO4 phase acts as a sintering inhibitor by separating Fe3O4particles and suppressing their sintering during oxidation. After reduction, the contact between the α-Fe and λ-Fe2W phases, forming ligament structures within each lamella, also hinders phase separation in the metallic state (FIG.28, (g)). Thus, W addition (>10 at%) enhances the stability of Fe lamellar structures effectively at two levels: (i) atomic-scale bonding of Fe and W, i.e., formation of λ-Fe2W and FeWO4 in the reduced and oxidized states respectively, to prevent W segregation; (ii) micro-scale mixture of α-Fe / λ-Fe2W in the reduced state and Fe3O4 / FeWO4 in the oxidized state to form ligaments and thus suppress phase separation and associated densification. In the early attempts to use Fe powder beds in ROBs, rapid sintering of Fe and Fe3O4phases at elevated temperatures (500‒800 ⁰C) resulted in a fast drop of the battery capacity to 65% after ten cycles. The present Fe-18W and Fe-25W foams, due to their enhanced resistance to sintering and fast reaction Attorney Docket No.: 616146.100555 kinetics during redox cycling at 800 °C, are promising for the applications as energy storage materials in ROBs for high-cycling energy storage with a long lifespan. Furthermore, Fe-W foams with stable porous structures during high-temperature redox cycling can be used in other gas-reforming technologies, e.g., chemical looping combustion, carbon utilization, and solar thermal redox. However, lamellar evolution is still observed after extensive cycling in both Fe-18W and Fe- 25W foams as a strong refinement of ligament sizes and as lamellar buckling. Specifically, in Fe-25W foams, the lamellar structure evolves into near-isotropic, highly porous, powder network, as shown in both cross-section and surface views (FIG.32) after 50 cycles. The CVT mechanism during the reduction and accumulation of vacancy during redox cycling increases microporosity within lamellae significantly, thus degrading the lamellar structure and leading to macroscopic foam swelling. This structural degradation is also reported in pure Fe, Fe-Ni, and Cu-Ni microparticles during high- temperature redox cycling. Mitigating such porosity generation to stabilize lamellar structures may lead to further improvements in long-term redox cycling capability. Conclusions Our study investigates the role of tungsten in enhancing the sintering resistance of Fe-W foams with hierarchical porous structures during redox cycling. The W concentration in freeze-cast, lamellar Fe-W foams was varied from 10 to 33 at% to achieve a broad range of Fe / Fe2W phase fractions and to investigate associated effects on the reaction performance and structural evolution of foams during redox H2 / H2O cycling at 800 ⁰C. The following conclusions are reached: -Increasing W concentration has a minor effect on the oxidation rate but significantly extends theinitial reduction time in the first cycle due to the slower reduction of more stable FeWO4, as compared to Fe3O4. -The chemical vapor transport (CVT) mechanism during reduction refines stable W-rich phases toa nanocrystalline scale, boosting the reaction kinetics of subsequent redox cycles. The reduction time could be adjusted by designing W concentrations and associated phase fractions of FeWO4. -Higher W concentration enhances the degradation resistance of Fe-W lamellar foams during high-temperature redox cycling, resulting in a more stable lamellae structure. oIn Fe-10W foam, sintering and densification of lamellar structures are notable via core-shell segregation after 20 redox cycles, similar to the degradation observed in Fe, Fe-Co, Fe-Cu, and Fe-Ni foams. During oxidation, Fe atoms diffuse outwards – microscopically to the surface of lamellae and macroscopically to the surface of the foams – which leads to the sintering of neighboring buckled lamellae and reduction in gas permeability. Attorney Docket No.: 616146.100555 oIn Fe-18W and Fe-25W foams, the W-rich phases increase in volume fraction and theyare in contact with the Fe-rich phases as (i) λ-Fe2W / α-Fe in the reduced state and (ii) FeWO4 / Fe3O4 in the oxidized state. These dual-phase microstructures in ligaments within lamellae prevent Fe out-diffusion and shell formation even after 50 redox cycles. oAlthough increasing W concentration from 10 to 18 at% enhances the stability of thelamellar structure effectively in Fe-W foams, lamellar buckling still occurs in Fe-18W foam, inducing swelling of foams after 50 cycles. -Our work establishes a compositional window for designing Fe-W foams with high degradationresistance for long-term redox cycling for energy storage. EXAMPLE 4: MICROSTRUCTURE AND PHASE EVOLUTION OF Fe-20Ni-20W FOAMS DURING HIGH- TEMPERATURE REDOX CYCLING This example discloses freeze-cast Fe-20Ni-20W (at%) foams suitable for high-temperature redox cycling showing excellent microstructural stability and regenerative porosity formation, but they exhibit slower reduction kinetics as compared to previously studied Fe-25W foams in EXAMPLE 1. The Fe-20Ni-20W lamellar foams, after initial hydrogen reduction, include a two-phase mixture of µ- Fe7W6 and γ-Fe(Ni,W), with significant microporosity due to the sintering inhibition of W. During oxidation by steam at 800 °C, these phases are oxidized to a three-phase mixture of (Fe,Ni)WO4, Fe3O4, and γ-Ni(Fe); upon subsequent reduction by H2, the foams return to their initial composition. The chemical vapor transport reduction of FeWO4results in the formation of submicron pores during each reduction half-cycle which accelerate subsequent reaction and limit sintering in the reduced state. The reduction of mixed oxide is relatively sluggish, which is likely due to the increased stability of the (Fe,Ni)WO4phase brought on by the substitution of Ni in Fe sites in FeWO4. Specifically, we investigate the high-temperature redox cycling behavior of Fe-20Ni-20W (at%) freeze-cast foams. This composition was chosen to ensure both high W content for sintering inhibition, and high Ni content to accelerate reduction. As described in detail below, we find that, while the foams retain the excellent redox stability expected from the high W content, the Ni addition does not accelerate reduction, while also slowing oxidation. This unexpected lack of effect from Ni is attributed to the formation of ternary Fe-W-Ni-based oxides with high stability. Fe-20Ni-20W (at.%) lamellar foams were fabricated by preparing an aqueous (87.5 vol% H2O) slurry composed of oxide powders – 5.11 vol% Fe2O3 (Noah Technologies, 99.9%, < 3 µm), 1.25 vol% NiO (Skyspring nano, 99.9%, 50 nm), and 3.63 vol% WO3 (Skyspring nano, 99.5%, < 100 nm) - in the Attorney Docket No.: 616146.100555 correct proportion to form the desired final composition after reduction; micrographs of the oxide powders.0.5 vol% Zephrym PD4974 (Croda) was used as a dispersant to prevent agglomeration of the oxide particles. The slurry was ball milled with yttrium-stabilized zirconia milling media for 24 hours to ensure good mixing of the powders, then combined with the polymer binder (2 vol% polyethylene glycol (Mn=3350, Sigma Aldrich) ), directionally freeze-cast, freeze-dried, debinded, and reduced under H2 according to EXAMPLE 1 above. Redox cycling was performed at 800 °C - under flowing H2O for oxidization, and under flowing H2for reduction - with details disclosed in EXAMPLE 1 above. In-situ X-ray diffraction (XRD) during redox cycling was performed on a small portion of a foam under conditions similar to the bulk redox cycling. Following a ramp at 30 °C / min to 800 °C, data were acquired in 1-minute spectra until the experiment was completed. Unlike previous studies which used Ar-4%H2as the reducing gas, pure H2was used for our in situ experiment, because of the increased thermodynamical stability of the oxidized state as discussed below. Microstructural characterization (SEM) used the same techniques as disclosed in EXAMPLE 1 above. In situ XRD results for Fe-20Ni-20W freeze cast foams are shown in FIG.33 for reduction at 800 ºC by pure H2 and oxidation by steam with Ar as a carrier gas. FIG.33, (a) shows time resolved XRD spectra for the first two redox cycles. The corresponding normalized peak integrals for the strongest peak of each phase are shown for the metallic phases (FIG.33, (b)) and for the oxide phases (FIG.33, (c)). The four blanks in the data are times when diffraction spectra were not measured while the gas environment was changed, corresponding to the change between a reducing environment (abbreviated red) and an oxidizing environment (abbreviated ox). The foam initially includes a two-phase mixture of μ-Fe7W6 and γ-Fe(Ni), after reduction at 1200 ºC from the freeze-cast oxides. As shown in FIG.33, (a)-(c), oxidation begins at t=5 min, immediately upon H2O (steam) exposure and proceeds rapidly, such that the formation of the transient FeO phase is not resolved at the scan time of 1 min. FeWO4 first forms 2 min into oxidation, with subsequent, slower formation of Fe3O4 at 5 min into oxidation. The formation of Fe3O4 is accompanied by a shift in the lattice parameter of the γ-Ni solid solution (to higher value, for the peak at 2θ=15.6º), as Fe in the matrix is oxidized and Ni becomes the majority element. The Fe-Ni-O phase diagram indicates that, at full oxidation, the γ-Ni(Fe) solid solution contains ~12% Fe. At t=19 min, the gases are switched and reduction under H2begins. Fe3O4is reduced first, in 6 min, with a corresponding return of the γ-Ni lattice parameter to its initial state. Reduction of FeWO4takes a much longer time (23 min), forming nano-crystalline Fe7W6, resulting in broad, poorly resolved peaks rather than the initially sharp peaks observed for t < 5 min (before the start of oxidation). These nanocrystals influence Attorney Docket No.: 616146.100555 successive cycles, as shown by the second oxidation and reduction: both are accelerated, with oxidation completed in 10 min, and reduction completed in 17 min. The transient re-formation of Fe3O4 during the reduction period (as seen by data at t=35, 65, and 70 min) indicates that there may be some gas flow tortuosity, resulting in evolved steam re-oxidizing previously reduced portions of the sample. The acceleration of reaction kinetics after the first redox cycle is consistent with previous studies of Fe-25W and Fe-25Mo lamellar foams. It is attributed to an increase in surface area and decrease in diffusion distance associated with the formation of nanocrystalline intermetallic phases and corresponding submicron porosity. While this is also observed for Fe-20Ni-20W, the overall reduction rate is much slower than expected for reduction with pure H2. For comparison, reduction of Fe-25W under pure H2takes 15 min for the first cycle, and 7 min for the second cycle, compared to present results (23 and 17 min for Fe-20Ni-20W, respectively). This result is unexpected, given the reduction acceleration effect as compared to pure Fe observed in the Fe-25Ni binary system. This slowdown in reduction is attributed to increased stability of the mixed oxide FeWO4. While a distinct FexNiyWOz compound is not formed, substitution of Ni into the Fe sites of FeWO4is likely (forming (Fe,Ni)WO4), and a similar substitution is observed in the intermetallic Fe7W6 (forming (Fe,Ni)7W6). This Ni substitution is expected to increase the stability of the mixed oxide via configurational entropy, and it appears to be sufficient to negate any gains in reduction kinetics compared to the binary Fe3O4reduction. Microstructurally, the Fe-20Ni-20W foams closely resemble previously studied Fe-25W foams. After freeze-casting and reduction of the blends of oxide powder followed by sintering at 1200 °C, the microstructure includes a fully densified, interpenetrating network of γ-Fe(Ni) and µ-Fe7W6. Based on the ternary Fe-Ni-W phase diagram, an equimolar mixture of γ-Fe(Ni,W) and µ-(Fe,Ni)7W6 is expected, and this phase mixture is observed in the as-fabricated foams. Both phases contain a high amount of Fe, the difference in the observed size is a result of the difference in sintering and coarsening behavior. The γ-Fe(Ni,W) contains relatively little W, and will thus coarsen and sinter much more than the intermetallic phase with high W content. Because of this, the γ-Fe(Ni,W) regions are larger. This initial state is shown in FIG.35 at three different magnifications. The overall foam architecture, shown from a radial cross-section in FIG.35, (a) includes metallic lamellae with uniform thickness separated by channels with varying width depending on the deformation of the neighboring lamellae. FIG.35, (b) shows a polished radial cross section of a group of lamellae (yellow letter L), each between 10-20 µm in thickness, separated by channels (letter C) templated by ice during the freeze casting step. The lamellae have significant internal porosity, as shown in FIG.35, (c). The lamellar microstructure includes an interpenetrating network of solid ligaments (red letter L) and micropores (letter M). The Attorney Docket No.: 616146.100555 solid ligaments are composed of µ-Fe7W6(white) and γ-Fe(Ni,W) (dark gray). Both phases are present in ligaments <10 µm in size. The high volume fraction of both phases indicates that they are interpenetrating. FIG.35, (d) shows a single, unmounted lamella after reduction and sintering. The same two-phase network is visible, and the presence of significant microporosity is observed as well, forming an interconnected 3-dimensional porous network. This porosity remains from the volume contraction associated with the reduction of the green body oxide powders, and it does not sinter due to the sintering inhibition of the µ-Fe7W6phase. Because of the sintering inhibition effects of W, the lamellae are initially porous, and they become increasingly porous with cycling. The overall utilization of Fe in each foam was determined by comparing the mass of the foam before cycling, after oxidation, and after reduction. All foams were able to fully oxidize and reduce in each 90-minute oxidation and 90-minute reduction period, up to the measured 50 cycles. Microstructural characterization did not reveal any directionality with respect to the foam reaction (e.g., residual oxidation at the top of the foam). The microstructure of a single lamella for foams subjected to 1, 20, and 50 cycles, in the reduced and oxidized states, is shown in FIG.36. The in situ XRD results (FIG.33) indicate the formation of nanocrystalline µ-Fe7W6 after the first oxidation / reduction cycle. This reduced state is shown in FIG.36, (a). The microporosity present within each lamella is significantly increased: the microporous network (orange arrows) expands in volume, and submicron pores (red arrows) are visible between metallic ligaments (red letter L) comprising the lamella. The micropores probably expand due to gas egress: on reduction, steam is produced at free surfaces around each open micropore with gas access and it escapes through the microporous network, causing the pores to expand; this process has been observed in other redox cycling studies as well. Because the µ-Fe7W6 phase is a strong sintering inhibitor, the newly-created and expanded micropores do not disappear by sintering after each cycle; this leads to a progressive increase in the porosity of the lamellae with continued cycling. The submicron pores form due to the reduction mechanism of FeWO4: this phase reduces in the presence of H2 by chemical vapor transport (CVT) forming a transient hydrated vapor phase (WO3(OH)2) which further reacts with H2 to condense as W on the surrounding Fe free surfaces with gas access, then reacts to form µ-Fe7W6. This gas- phase-precipitated µ phase is nanocrystalline, as indicated by the XRD results, with submicron pores separating neighboring nanocrystalline µ-phase grains. Like the micropores, these submicron pores do not sinter due to the high sintering inhibition of the µ phase that borders them. In the oxidized state after 1 cycle (FIG.36, (b)), the microporous network remains present (orange arrows), though it is somewhat filled because of the volumetric as the metallic phase oxidizes to FeWO4 and Fe3O4. Formation of FeWO4 whiskers (yellow arrows) is observed as well. Attorney Docket No.: 616146.100555 After 20 cycles in the reduced state (FIG.36, (c)), the microporous network has expanded further, with each lamella more closely resembling a fully porous powder foam rather than a sintered platelet. The cyclic redistribution of W during each reduction cycle shrinks the size of the γ and µ phases, with each becoming approximately micron-sized. Submicron pores are prevalent throughout these micron size phases. The oxidized state (FIG.36, (d)) closely resembled the reduced state, with individual Fe3O4 and FeWO4 particles observed on the lamellar surface. After 50 cycles, the microstructure has not changed, with similar pore and particle size in the reduced (FIG.36, (e)) and oxidized (FIG.36, (f)) states. Throughout cycling, the metallic Ni(Fe) phase which does not oxidize remains well-distributed throughout the microstructure. In conclusion, Fe-20Ni-20W lamellar foams are resistant against degradation upon long-term redox cycling, with no sintering or elemental segregation observed after 50 cycles. The lamellae, which are not fully dense initially, grow progressively more porous with cycling due to the CVT mechanism, and the size of each phase decreases from ~5 µm in the as-fabricated state to micron or submicron after 20 cycles, with no significant change observed between 20 and 50 cycles. The metallic Ni phase does not provide a strong accelerating effect on the reduction process, making its addition unnecessary in view of similar microstructural results obtained for binary Fe-25W lamellar foams fabricated via the same freeze-casting process. The lack of improvement in reduction kinetics can be explained by (i) the CVT reduction mechanism of FeWO4, which does not benefit from the presence of Ni, and (ii) the formation of more stable quaternary (Fe,Ni)WO4 oxides due to the presence of Ni, which are slower to reduce than ternary FeWO4. Thus, while Fe3O4reduces rapidly, the overall rate of reduction of the entire phase system does not benefit from the presence of Ni, due to the formation of this more stable oxide phase. The foregoing description of the exemplary embodiments of the invention has been presented only for the purposes of illustration and description and is not intended to be exhaustive or to limit the invention to the precise forms disclosed. Many modifications and variations are possible in light of the above teaching. The embodiments were chosen and described to explain the principles of the invention and their practical application to enable others skilled in the art to utilize the invention and various embodiments and with various modifications as are suited to the particular use contemplated. Alternative embodiments will become apparent to those skilled in the art to which the invention pertains without departing from its spirit and scope. Accordingly, the scope of the invention is defined by the appended claims rather than the foregoing description and the exemplary embodiments described therein.

Claims

Attorney Docket No.: 616146.100555 CLAIMS What is claimed is:

1. A redox system, comprising:a Fe-xW powder architecture fabricated to possess excellent resistance to degradation at a high-temperature during redox cycling between oxidized and reduced states, wherein x represents a concentration (at%) of tungsten (W) in the powder architecture and is in a range of about 1-75 at%, wherein the powder architecture comprises powder bed, powder suspension, foams, fibers, and / or printed microlattice.

2. The redox system of claim 1, wherein the redox cycling is performed with oxidizing gases andreducing gases alternatively, wherein the oxidizing gases are those that increase the oxidation state of iron (Fe) and / or tungsten (W) and comprises H2O, CO2, and / or O2, and wherein the reducing gases are those that decrease the oxidation state of Fe and / or W and comprises H2, hydrocarbons including CH4, ethane, propane and / or butane, NH3, and / or CO.

3. The redox system of claim 2 wherein the oxidizing gas and the reducing gas are combined withany inert gases.

4. The redox system of claim 1, wherein the high-temperature is in a range of about 500 – 1500°C.

5. The redox system of claim 1, wherein the powder architecture possesses the excellent resistanceto degradation at about 800 °C during the redox cycling between the metallic and oxide states, with good reaction kinetics maintained up to a minimum of 100 full cycles.

6. The redox system of claim 5, wherein the powder architecture is of Fe-1W, Fe-5W, Fe-10W,Fe-18W, Fe-20W, Fe-25W, Fe-30W, Fe-33W, Fe-40W, Fe-45W, Fe-50W, or Fe-75W.

7. The redox system of claim 6, wherein the Fe-10W porous powder architecture with a lowervolume fraction of W-containing phases of λ-Fe2W and FeWO4has lamellae densification andAttorney Docket No.: 616146.100555 core-shell structure formation, due to Fe outward diffusion during oxidation.

8. The redox system of claim 6, wherein each of the Fe-18W and Fe-25W powder architecturescomprises a sufficient volume fraction of W-rich phases such as λ-Fe2W to inhibit sintering for α-Fe in the reduced state and FeWO4to inhibit sintering for Fe3O4in the oxidized state, thereby forming a porous solid comprising two phases of Fe / λ-Fe2W and Fe3O4 / FeWO4, respectively.

9. The redox system of claim 8, wherein each porous region in the Fe-25W powder architectureinitially includes a mixture of bcc α-Fe(W) and µ-Fe7W6; wherein during oxidation, α-Fe oxidizes to Fe3O4 and µ-Fe7W6 oxidizes to FeWO4 and minor amounts of Fe3O4; and wherein during reduction, the Fe3O4 reduces back to bcc Fe, and the FeWO4 reduces more slowly to Fe2W.

10. The redox system of claim 9, wherein there is a mismatch in the stoichiometric ratio between Feand W in moving from FeWO4(Fe:W=1) to Fe2W (Fe:W=2), whereby Fe from the surrounding regions from prior reduction of Fe3O4is used to form the intermetallic Fe2W.

11. The redox system of claim 9, wherein the second oxidation is, on average, about 40-80% fasterthan the first oxidation, and the second reduction is, on average, about 10-50% faster than the first reduction.

12. The redox system of claim 9, wherein in the reduced state, the powder architecture includespowders which exhibit 2-10 µm ligaments interpenetrated by open porosity of the same size, wherein each ligament, in turn, includes an α-Fe(-1.6W) backbone decorated, and at times interrupted, with submicron (1-900 nm) λ-Fe2W grains and submicron pores formed by chemical vapor transport, and wherein this same reduced microstructure is maintained through 100 cycles with the micron-scale open porosity increasing with cycle number.

13. The redox system of claim 9, wherein the oxidized state includes two interpenetrating networksof Fe3O4 and FeWO4.Attorney Docket No.: 616146.10055514. The redox system of claim 1, wherein the powder architecture, after the first cycle, comprisesthree levels of pore sizes including macroscopic freeze-cast channels (level I) separating lamellae which are distributed into colonies with aligned radial orientation; microscale sintering-inhibition pore network (level II) within each lamella, fully continuous, formed due to the sintering inhibition of W; and submicron chemical vapor transport (CVT) pores (level III), formed within each W-containing area in the lamellae, regenerated at each reduction half-cycle, thereby creating a hierarchical porosity structure in the powder architecture.

15. The redox system of claim 14, wherein the microscale sintering-inhibition pores increase thesurface area and thus reactivity of the ligaments and the submicron CVT pores greatly increase surface area and limit damage accumulation by effectively regenerating the original microstructure after each cycle, producing nanocrystalline λ–Fe2W with corresponding nanopores.

16. The redox system of claim 15, wherein the microstructural stability allows for the powderarchitecture to grow progressively more porous (rather than denser) with cycling, and the W- rich phases become both smaller in morphology and more homogeneously distributed throughout the ligaments, thereby increasing the sintering inhibition effect and generating more submicron pores to accelerate subsequent reaction.

17. The redox system of claim 14, wherein the resistance is attributed to the sintering inhibition ofW combined with the CVT mechanism of reduction.

18. The redox system of claim 14, wherein the resistance is attributed to both the sinteringinhibition inherent to W, and the nature of the intermetallic and mixed oxide compounds formed between Fe and W, such that W does not segregate and remains mixed with Fe at the micron or submicron scale, at all times.

19. The redox system of claim 18, wherein W is mixed with Fe at the micron or submicron scale,either in the intermetallic λ–Fe2W or in the mixed oxide FeWO4 phases, thereby effectively preventing W from segregating into a pure W phase with micron to millimeter size, that reducesAttorney Docket No.: 616146.100555 the degradation resistance of the Fe-25W powder architecture.

20. The redox system of claim 1, wherein increasing W concentration has a minor effect on theoxidation rate but significantly extends the initial reduction time in the first cycle due to the slower reduction of more stable FeWO4, as compared to Fe3O4.

21. The redox system of claim 1, wherein the chemical vapor transport (CVT) mechanism duringreduction refines stable W-rich phases to a nanocrystalline scale, boosting the reaction kinetics of subsequent redox cycles. The reduction time can be adjusted by designing W concentrations and associated phase fractions of FeWO4.

22. The redox system of claim 1, wherein higher W concentration enhances the degradationresistance of Fe-W lamellar foams during high-temperature redox cycling, resulting in a more stable structure.

23. The redox system of claim 1, wherein the powder architecture further comprises Ni, Co, Cu,and / or Mo, in compositional ranges Fe-(5-33)(Mo+W), Fe-(5-33)(Ni+W), or Fe-(5- 33)(Mo+Ni+W), all in at%.

24. The redox system of claim 23, wherein the powder architecture is of Fe-(5-33)(Mo, W),wherein Mo and W are both active in oxidation and reduction.

25. The redox system of claim 24, wherein the powder architecture is of Fe-9Mo-9W.

26. The redox system of claim 23, wherein the powder architecture is of Fe-(5-33)(Ni, W), whereinW is active in oxidation and reduction while Ni is inert.

27. The redox system of claim 26, wherein the powder architecture is of Fe-20Ni-20W.

28. The redox system of claim 26, wherein the Fe-20Ni-20W powder architecture, after initialreduction, include a two-phase mixture of µ-Fe7W6 and γ-Fe(Ni,W), with significantAttorney Docket No.: 616146.100555 microporosity due to the sintering inhibition of W, and wherein during oxidation at 800 °C, these phases are oxidized to a three-phase mixture of (Fe,Ni)WO4, Fe3O4, and γ-Ni(Fe); upon subsequent reduction by H2, the powder architecture returns to their initial composition.

29. A device, comprising the redox system of any one of claims 1-28.

30. A method for fabricating a redox system, comprising:milling a blend of an iron (Fe) or Fe-containing powder and a tungsten (W) or W- containing powder into a homogenous powder; and optionally heat treating the powder blend to yield a powder architecture of Fe-xW, wherein x represents a concentration (at%) of W in the powder architecture and is in a range of about 1-75 at%.

31. The method of claim 30, wherein the Fe or Fe-containing powder comprises a Fe powder, a Fe-rich powder, a Fe-containing oxide powder, a Fe-containing nitride powder, and / or a Fe- containing carbide powder, and wherein the W or W-containing powder comprises a W powder, a W-rich powder, a W-containing oxide powder, a W-containing nitride powder, a W- containing carbide powder, and / or ammonium paratungstate.

32. The method of claim 30, wherein the powder architecture is of Fe-1W, Fe-5W, Fe-10W, Fe-18W, Fe-20W, Fe-25W, Fe-30W, Fe-33W, Fe-40W, Fe-45W, Fe-50W, or Fe-75W.

33. The method of claim 30, wherein the powder architecture possesses the excellent resistance todegradation at about 800 °C during redox cycling between oxidized and reduced states, with good reaction kinetics maintained up to a minimum of 100 full cycles.

34. A method for fabricating a redox system, comprising:providing a Fe-xW alloy; atomizing the alloy to form a powder; and optionally heat treating the powder to yield a powder architecture of Fe-xW, wherein x represents a concentration (at%) of W in the powder architecture and is in a range of about 1-75Attorney Docket No.: 616146.100555 at%.

35. The method of claim 34, wherein said providing the Fe-xW alloy is performed by melting ironand ferrotungsten.

36. A method for fabricating a redox system, comprising:producing a Fe- and W-rich oxide by sintering of precursor powders; pulverizing the Fe- and W-rich oxide to form a powder with Fe-xW oxide composition; and optionally heat treating the powder to yield a powder architecture with metallic Fe-xW composition, wherein x represents a concentration (at%) of W in the powder architecture and is in a range of about 1-75 at%.

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