Reduced early transition metal oxides/metal composites have been identified here as interesting thermoelectric materials. Numerous compositions in the Nb-rich portion of the WO3–Nb2O5 system have been studied, in composite formulations with elemental W. Spark plasma sintering (SPS) has been employed to achieve rapid preparation and consolidation of composite materials containing W metal precipitates with characteristic length scales that range from under 20 nm to a few microns, that exhibit thermal conductivities that are constant from 300 K to 1000 K, approximately 2.5 W m−1 K−1. Thermoelectric properties of these n-type materials were measured, and the highest-performing compositions were found to reach figure of merit zT values close to 0.1 at 950 K. The measurements point to higher zT values at yet-higher temperatures.

The promise of solid-state energy conversion and the recovery of energy from waste heat in devices that have no moving parts has led to vigorous research on thermoelectric materials. However, widespread deployment of thermoelectric modules has been limited by some combination of low efficiency, low crustal abundance of the key elements in routinely-studied thermoelectrics, and poor thermal stability of materials. Oxide materials potentially help address some of these issues although reduced metal oxide materials require oxidation protection to maintain desirable electrical properties. Despite the opportunity for oxide thermoelectrics, few high-performance n-type materials have been discovered. The performance of thermoelectric materials is usually measured by the figure of merit zT given by zT = (S2T)/(ρκ); a function of the electrical resistivity ρ, the Seebeck coefficient S, and the thermal conductivity κ of the material at temperature T. A recent datamining study carried out by some of us,1 had pointed out that empirically, high performing thermoelectric materials operating at (typically) 1000 K have Seebeck coefficients in the range of 200 μVK−1 to 300 μVK−1, and electrical resistivities close to 1 × 10−3 Ωcm. In many oxide materials, attaining high Seebeck coefficients is frequently not as challenging as achieving these relatively metallic resistivity values.2 In addition, small-unit-cell oxides such as spinels, perovskites, and rutiles, are usually highly connected, and therefore have relatively high lattice thermal conductivities, which is undesirable. In this study, early transition-metal oxides with relatively large unit cells, prepared as composites with metals, have been rationally selected as potentially being able to overcome the pitfalls of high resistivity coupled with low thermal conductivity that many bulk oxides display.

The WO3–Nb2O5 phases selected for study here tend to form complex Wadsley–Roth shear structures, which feature 3-dimensional slabs of corner-sharing NbO6 octahedra connected by WO4 tetrahedra.3 Nb2O5 has many polymorphs, but the dominant room-temperature phase is a layered structure with distorted and irregular Nb polyhedra.4 This poor structural connectivity is not ideal for electrical transport and electronic doping. However, WO3–Nb2O5 phases can be doped to have low electrical resistivity, in part because of the highly connected 3D structures that derive from the perovskite aristotype. Highly connected structures also tend to have high thermal conductivity, but this is mitigated in the Wadsley–Roth structure-types because of the complex, large-unit cell structures. In the study here, with W/Nb substitution, oxygen deficiency, and crystallographic shear (Figure 1) also potentially decrease thermal conductivity. Indeed, Winter and Clarke have shown the thermal conductivity of stoichiometric d0 analogues to the materials studied here to be about 2 W m−1 K−1 from 300 K to 1300 K.5 

FIG. 1.

WO3–Nb2O5 phases tend to form complex Wadsley–Roth shear structures, which feature 3-dimensional slabs of corner-sharing NbO6 octahedra connected by WO4 tetrahedra. W3Nb14O44 is one such example of many possible homologues; the size of the slabs can be changed with W/Nb composition. The high polyhedral connectivity is important for good electrical transport, while the intrinsic structural defects (such as the crystallographic shear planes and W/Nb site substitution) are important for low thermal transport.

FIG. 1.

WO3–Nb2O5 phases tend to form complex Wadsley–Roth shear structures, which feature 3-dimensional slabs of corner-sharing NbO6 octahedra connected by WO4 tetrahedra. W3Nb14O44 is one such example of many possible homologues; the size of the slabs can be changed with W/Nb composition. The high polyhedral connectivity is important for good electrical transport, while the intrinsic structural defects (such as the crystallographic shear planes and W/Nb site substitution) are important for low thermal transport.

Close modal

In this contribution, the thermoelectric properties of composite samples from the Nb-rich portion of the WO3–Nb2O5 system are reported. The samples were prepared in single step of short duration that combined preparation from starting materials (elemental metals and oxides) with consolidation. The technique employed for the one-step preparation and consolidation was spark-plasma sintering or SPS, perhaps more accurately referred to as current-assisted pressure-activated densification.6 This is a rapid preparatory route to react materials, with typical reaction taking minutes, instead of hours or days, and yielding dense, consolidated solids that can be directly employed for physical property measurements. The preparation of phase-pure single-cation WO3−δ7 and multi-cation Cr2WO68 have been previously demonstrated using SPS, proving that ion mobility is sufficient to allow diffusion and phase equilibrium in WO3 systems at these reaction temperatures and time-scales. Rapid preparation by SPS is favourable to conventional solid-state preparation of the W–Nb–O oxide materials studied here, which conventionally require extended heat treatments at high temperatures (e.g., 1350°C in sealed platinum crucibles).9 In addition to being time-consuming, extended heat treatments at high temperatures also result in difficulties with stoichiometry control because of volatile oxide precursors; a problem often encountered in the preparation of oxides of W.9 Furthermore, the final product of conventional preparation is usually powder that must be consolidated and densified before physical property measurements can be performed. In this work, SPS was used to prepare dense pellets of (W1.083O3)x(Nb2.1O5)1−x from a mixture of binary oxides and metal powders in less than 30 min per sample.

Fourteen compositions were chosen based on reported stable phases,10 even though the crystal structure was not known in all cases. Stoichiometric WO3, Nb2O5, or (WO3)x(Nb2O5)1−x (d0 oxides) are electrical insulators, so the materials were doped with 0.5 electrons per metal cation by adding corresponding amounts of elemental metal to the initial mixture. Tungsten was treated as a d6 donor, so 1/12 molar equivalents of W were added for every equivalent of WO3—1/10 equivalents of Nb, treated as d5, were added for each Nb2O5. The final compositions of (W1.083O3)x(Nb2.1O5)1−x are consistent with the W/Nb ratio determined after preparation, by elemental analysis using energy-dispersive X-ray spectroscopy (EDX).

Polycrystalline samples were made by simultaneous reaction and consolidation of binary oxide and elemental powders (WO3, 99.9%, Sigma-Aldrich; Nb2O5, 99.99%, Sigma-Aldrich, W, 99.99%, Sigma-Aldrich; Nb, 99.8%, Sigma Aldrich) using spark plasma sintering on an instrument from FCT Systeme GmbH, Germany. In a chamber base pressure of 10 torr of argon, 80 MPa uniaxial pressure was applied using a 9 mm inner-diameter graphite die (EDM-4, POCO graphite). The sample was then heated rapidly to 1473 K at 150 K/min, annealed for 6 min at 1473 K, then rapidly cooled at 180 K/min (Figure 2). X-ray diffraction (XRD) of the resulting dense pellets was performed using a laboratory instrument (Philips X’Pert diffractometer, Cu Kα radiation). After initial room-temperature electrical transport measurements, a subset of these samples was subjected to more rigorous study (Figure 3).

FIG. 2.

A typical SPS densification profile during sample processing. SPS was used to prepare dense pellets of (W1.083O3)x(Nb2.1O5)1−x composites from a mixture of binary oxides and metal powders in under 30 min. 80 MPa of uniaxial pressure was applied and the sample was heated rapidly from 673 K to 1473 K at 170 K/min, annealed for 6 min at 1473 K, then rapidly cooled at 180 K/min. The temperature profile below 673 K was estimated, as the pyrometer is unreliable below 673 K.

FIG. 2.

A typical SPS densification profile during sample processing. SPS was used to prepare dense pellets of (W1.083O3)x(Nb2.1O5)1−x composites from a mixture of binary oxides and metal powders in under 30 min. 80 MPa of uniaxial pressure was applied and the sample was heated rapidly from 673 K to 1473 K at 170 K/min, annealed for 6 min at 1473 K, then rapidly cooled at 180 K/min. The temperature profile below 673 K was estimated, as the pyrometer is unreliable below 673 K.

Close modal
FIG. 3.

SPS was used to prepare 14 samples that were subsequently characterized to elucidate the evolution of thermoelectric properties with composition in the (W1.083O3)x(Nb2.1O5)1−x phase system. After measurement of room temperature electrical properties, smaller subsets of samples were investigated more thoroughly.

FIG. 3.

SPS was used to prepare 14 samples that were subsequently characterized to elucidate the evolution of thermoelectric properties with composition in the (W1.083O3)x(Nb2.1O5)1−x phase system. After measurement of room temperature electrical properties, smaller subsets of samples were investigated more thoroughly.

Close modal

High-resolution synchrotron XRD data on finely-ground powder was acquired at 295 K at beamline 11-BM at the Advanced Photon Source (APS), Argonne National Laboratory, using an average wavelength of ∼0.4137 Å with a diffractometer that has been described in detail by Wang et al.11 The precise wavelength was determined using a mixture of Si (SRM 640c) and Al2O3 (SRM 676) NIST standards run in a separate measurement. Samples were contained in 0.8 mm diameter Kapton capillaries and the packing density was low enough that absorption was not noticeable. Rietveld refinement was performed using the TOPAS 5 Academic program suite.12 Instrument profile parameters were determined using Si (SRM 640c) and LaB6 (SRM 660b) NIST lineshape standards, run in separate measurements.

Scanning electron microscopy (SEM) was performed on an FEI XL30 microscope using a backscatter electron (BSE) detector; EDX data were collected with an EDAX Si-drift detector. Typical accelerating voltages were 15 keV or 20 keV. To prepare samples for SEM, pellets were mounted in epoxy and polished with successively finer abrasives; the final polishing step was performed with 0.25 μm colloidal diamond suspension on MicroCloth. A thin layer of Pd–Au, 2 nm to 5 nm, was sputtered on the surface to alleviate charging during imaging. The brightness and contrast of the micrographs presented have not been post-processed in any way. Particle size analysis was performed using ImageJ13 on 2500 × magnification micrographs, areas of approximately 33 μm × 33 μm, to visualize the smallest features, some of which remain sub-pixel diameter even at this magnification. A minimum size cutoff threshold of 0.05 μm2 was used to minimize false positives.

The high-temperature (950 K max) electrical resistivity and Seebeck coefficient were measured with an ULVAC Technologies ZEM-3. Samples had approximate dimensions of 9 mm × 3 mm × 3 mm. Measurements were performed in a helium under-pressure, and data was collected through three heating and cooling cycles to ensure sample stability and reproducibility. No changes in electrical properties were observed between cycles, and analysis of the lab XRD patterns from materials after measurements showed no changes in structure nor new phases. Additionally, comparing BSE micrographs collected before and after electrical property measurements revealed no change in microstructure.

The thermal conductivity, κ, was calculated from the density ρ, thermal diffusivity α, and heat capacity Cp, using the relation κ = ραCp. The densities of consolidated samples were measured using helium pycnometry (MicroMeritics AccuPyc 1340 Pycnometer). To reduce the large systematic errors often introduced by measuring the heat capacity,14 the temperature-dependent heat capacity was estimated using the empirical Neumann-Kopp rule based on the chemical composition and literature values for the heat capacity of WO3 and Nb2O5,15 as outlined by Leitner et al.16 The estimated error for the heat capacity introduced by this method was <3% at 300 K, and <5% at 1000 K. The thermal diffusivity (α) was measured using a Netzsch LFA 457 laser flash system. The samples for the thermal diffusivity measurements were machined to be coplanar with a thickness of between 1 mm and 2 mm and a diameter of 8 mm. The samples were then spray coated with colloidal graphite on both sides to ensure maximum optical absorption and emissivity. Thermal diffusivity was measured in an argon gas atmosphere from 308 K to 1050 K at intervals of 50 K. A final measurement at 308 K was taken after cooling to ensure reproducibility of the measurements and to ensure the coatings remained intact. The thermal diffusivity values were determined using the model of Clark and Taylor,17 which corrects for radiative losses.

Powder XRD reveals the presence of elemental, body-centered cubic (BCC) W in many samples, in addition to peaks characteristic of W–Nb–O phases (Figure 4). A sample refinement is shown in Figure 5, using monoclinic WNb12O33 (space group C2, with refined parameters very near those reported by Roth et al.9) and BCC W for the model. The fit is able to account for the elemental W but is unable to describe the primary phase reflections, particularly at low Q. The exact W–Nb–O phases were not able to be determined using powder XRD, as their crystal structures have large unit cells with many atomic positions, and structural similarity leads to many coexisting phases that complicate analysis. Furthermore, the XRD patterns of the many phases can only be distinguished by weak superstructure peaks at very low angles (Q < 1.5 Å−1) created by O ordering, which contributes only weakly to the total X-ray scattering compared to the more electron-rich W and Nb atoms in the materials.

FIG. 4.

Synchrotron XRD was performed on a subset of 8 samples in the (W1.083O3)x(Nb2.1O5)1−x series. Although specific W–Nb–O phases could not be determined due to the insensitivity of XRD to the many of the possible superstructures in these materials, the presence of BCC W in most samples is evident. The intensity of reflections from W metal increases with increasing x from x = 0.02 to x = 0.13.

FIG. 4.

Synchrotron XRD was performed on a subset of 8 samples in the (W1.083O3)x(Nb2.1O5)1−x series. Although specific W–Nb–O phases could not be determined due to the insensitivity of XRD to the many of the possible superstructures in these materials, the presence of BCC W in most samples is evident. The intensity of reflections from W metal increases with increasing x from x = 0.02 to x = 0.13.

Close modal
FIG. 5.

X-ray diffractogram and two-phase Pawley refinement of (W1.083O3)0.02(Nb2.1O5)0.98, modeled using WNb12O33 and W. Reflections from elemental W are well described, but reflections from the primary phase are poorly fit at low Q when using reported W–Nb–O crystal structures as starting models.

FIG. 5.

X-ray diffractogram and two-phase Pawley refinement of (W1.083O3)0.02(Nb2.1O5)0.98, modeled using WNb12O33 and W. Reflections from elemental W are well described, but reflections from the primary phase are poorly fit at low Q when using reported W–Nb–O crystal structures as starting models.

Close modal

SEM was performed to investigate the microstructure of the samples, and to reveal the disposition of the multiple phases in each sample. There were no noticeable differences between cross-sections taken from the top or sides of the pellets, suggesting the microstructures are uniform. In previous studies SPS has been used to prepare single-phase W-containing oxides in a single step.7,8 In contrast, when we employ a single step SPS process in this W–Nb–O system we observe multiple phases. BSE micrographs combined with EDX reveal the presence of W-rich inclusions with a wide range of length scales (Figure 6). Owing to the small precipitate size and large interaction volume of EDX, approximately 1 μm in diameter, EDX is unable to unambiguously determine the precise atomic composition, though it is likely these precipitates are W metal and give rise to the W metal reflections observed by XRD. The smallest inclusions are <20 nm, whereas the larger inclusions are approximately 500 nm along their narrow axis and up to 2 μm long. The dimensions of these precipitates are ideal for scattering phonons with a broad wavelength distribution and lead to low thermal conductivity at low temperature, as shown later.

FIG. 6.

Backscatter electron micrographs (W1.083O3)x(Nb2.1O5)1−x (x = 0.02, 0.13, 0.31) at 200 × and 2500 × magnification reveal small W metal precipitates (white) homogeneously distributed throughout samples with low W content (x ≤ 0.13). The smallest inclusions are sub-pixel diameter (<20 nm), whereas the larger inclusions are a few microns in length, covering a wide range of relevant length scales to interact with phonons. The average size and number of precipitates increases with x in this regime. This microstructure is stable after repeated thermal cycling up to 950 K, the highest temperatures measured. At high W content (x ≥ 0.31), the microstructure is poorly controlled, and there are larger domains of multiple minor phases, tens of microns in width).

FIG. 6.

Backscatter electron micrographs (W1.083O3)x(Nb2.1O5)1−x (x = 0.02, 0.13, 0.31) at 200 × and 2500 × magnification reveal small W metal precipitates (white) homogeneously distributed throughout samples with low W content (x ≤ 0.13). The smallest inclusions are sub-pixel diameter (<20 nm), whereas the larger inclusions are a few microns in length, covering a wide range of relevant length scales to interact with phonons. The average size and number of precipitates increases with x in this regime. This microstructure is stable after repeated thermal cycling up to 950 K, the highest temperatures measured. At high W content (x ≥ 0.31), the microstructure is poorly controlled, and there are larger domains of multiple minor phases, tens of microns in width).

Close modal

At low W content (x ≤ 0.13), W metal secondary phase inclusions are homogeneously distributed and quite small, with median sizes of a few hundred nanometers, suggesting possible solubility at high temperature and phase segregation due to precipitation upon rapid cooling (180 K/min) after the SPS hold at 1473 K. The average size and area coverage of the precipitates increases in a controlled fashion with W content (Figure 7). This is consistent with the increasing intensity of reflections from W metal seen by XRD in this range. Although precipitates of sub-pixel diameter (<20 nm) can be seen in micrographs, only precipitates larger than 0.05 μm2 (effective diameter of 250 nm) were used for this analysis to minimize false detection events. The precipitate sizes appear to follow log-normal distributions, though meaningful fitting to extract statistical descriptors would require more statistics of particles with smaller sizes.

FIG. 7.

The average size of W metal precipitates measured in BSE micrographs and the area fraction of precipitates increases smoothly with W content when x ≤ 0.13 in (W1.083O3)x(Nb2.1O5)1−x. To minimize false detection events, only precipitates larger than 0.05 μm2 (effective diameter of 250 nm) were used for this analysis.

FIG. 7.

The average size of W metal precipitates measured in BSE micrographs and the area fraction of precipitates increases smoothly with W content when x ≤ 0.13 in (W1.083O3)x(Nb2.1O5)1−x. To minimize false detection events, only precipitates larger than 0.05 μm2 (effective diameter of 250 nm) were used for this analysis.

Close modal

At high W content (x ≥ 0.31), in contrast, the microstructure is less controlled, with multiple minor phases, existing as larger domains in excess of 10 μm in width. The presence of these minor phases not observed at lower W loadings suggests that that these high-W compositions may sit in a three-phase field of the W–Nb–O ternary, while low-W compositions are two-phase. The W-rich inclusions transition from evenly dispersed sub-micron particles to unsystematic domains tens of microns in size. This is likely because the concentration and domain size of these inclusions increases with W content, up to some point at which they coalesce and segregate to minimize interfacial area, and therefore surface energy. This change in behavior may be related to the critical percolation threshold (27% site filling for a random close-packed three-dimensional lattice).18 

To determine the thermoelectric performance of the composites, electrical transport properties of (W1.083O3)x(Nb2.1O5)1−x were measured up to 950 K (Figure 8). The thermoelectric performance is defined by the thermoelectric figure of merit, zT = S2ρ−1κ−1T, which is a function of the material electrical resistivity ρ, Seebeck coefficient S, thermal conductivity κ, and temperature T. Through the compositional series, ρ increases with W content until x = 0.13, after which it decreases continually, up to the highest W content measured. S changes predictably in opposition to the electrical resistivity with composition. The power factor, S2ρ−1, increases monotonically with temperature up to at least 950 K, the highest temperature measured. With regard to sample composition, the maximum power factor is obtained for samples with x = 0.08.

FIG. 8.

(W1.083O3)x(Nb2.1O5)1−x displays metallic electrical transport when x ≤ = 0.08, and maximum electrical performance (S2/ρ) for thermoelectric considerations is seen at x = 0.08. The electrical transport properties of the materials do not change after repeated thermal cycling up to the maximum temperatures measured, 950 K. One-step SPS prepared x = 0.02 material (squares) has comparable properties to material of the same composition made using traditional solid-state methods (red line).

FIG. 8.

(W1.083O3)x(Nb2.1O5)1−x displays metallic electrical transport when x ≤ = 0.08, and maximum electrical performance (S2/ρ) for thermoelectric considerations is seen at x = 0.08. The electrical transport properties of the materials do not change after repeated thermal cycling up to the maximum temperatures measured, 950 K. One-step SPS prepared x = 0.02 material (squares) has comparable properties to material of the same composition made using traditional solid-state methods (red line).

Close modal

Several samples were also made by conventional solid-state reaction of precursors in evacuated silica ampoules at 937 K for 48 hours,10 followed by densification using the same SPS processing. Samples made using this two-step process displayed similar X-ray diffraction patterns and high-temperature electrical transport properties (electrical resistivity and Seebeck coefficient) as analogous samples made using a one-step process.

Thermal conductivity was determined for a smaller subset of samples (Figure 9) to determine the influence of the microstructures observed in BSE micrographs (Figure 6). Densities used to calculate thermal conductivity were obtained through He pycnometry on consolidated pellets. Although the relative density (compared to the theoretical maximum single-crystal density) cannot be unambiguously known due to the presence of multiple phases and oxygen deficiency, the samples examined here are highly dense; scanning electron microscopy of polished cross-sections reveals no evidence of significant porosity, and comparison with known phases of similar composition shows the composites are of similar density (Table I). The density of the composites is consistently higher by approximately 5%, owing to the reduced nature of the samples.

FIG. 9.

Very little W content, x = 0.02, is sufficient to decrease the thermal conductivity compared to Nb2O5, particularly at low temperatures where longer-wavelength phonons are the dominant carriers of heat. Despite the increase in electrical thermal conductivity, the total thermal conductivity x = 0.02 composites is reduced by 30% at 300 K compared to insulating Nb2O5. The thermal conductivity of these biphasic materials is largely invariant with temperature, likely because long-wavelength phonons at low temperatures are scattered by the W metal precipitates that span length scales of <20 nm to a few microns. At high W content, the same reduction in thermal conductivity is not seen.

FIG. 9.

Very little W content, x = 0.02, is sufficient to decrease the thermal conductivity compared to Nb2O5, particularly at low temperatures where longer-wavelength phonons are the dominant carriers of heat. Despite the increase in electrical thermal conductivity, the total thermal conductivity x = 0.02 composites is reduced by 30% at 300 K compared to insulating Nb2O5. The thermal conductivity of these biphasic materials is largely invariant with temperature, likely because long-wavelength phonons at low temperatures are scattered by the W metal precipitates that span length scales of <20 nm to a few microns. At high W content, the same reduction in thermal conductivity is not seen.

Close modal
TABLE I.

The densities of (W1.083O3)x(Nb2.1O5)1−x composites measured by He pycnometry are compared with the single crystal densities of known crystal structures with similar composition. The densities of composites are consistently high due to the reduced nature of the composites creating high density metal particles.

x nearest known crystal structure ρmeasured(g/cm3) ρcrystal(g/cm3) % density
Nb2O5  4.992(2)  4.55  110% 
0.02  Nb2O5  4.721(4)  4.55  104% 
0.08  WNb12O33  4.928(2)  4.76  104% 
0.13  W4Nb26O77  5.114(2)  4.93  104% 
0.53  W9Nb8O47  6.360(2)  5.93  107% 
x nearest known crystal structure ρmeasured(g/cm3) ρcrystal(g/cm3) % density
Nb2O5  4.992(2)  4.55  110% 
0.02  Nb2O5  4.721(4)  4.55  104% 
0.08  WNb12O33  4.928(2)  4.76  104% 
0.13  W4Nb26O77  5.114(2)  4.93  104% 
0.53  W9Nb8O47  6.360(2)  5.93  107% 

For (W1.083O3)x(Nb2.1O5)1−x samples with low W content (x ≤ 0.13), the thermal conductivity is approximately 2.5Wm−1 K−1, from 300 K to 1050 K (Figure 9). For comparison, the thermal conductivity of stoichiometric Nb2O5, prepared using the same SPS processing conditions, was measured as a reference. Although the thermal conductivity of Nb2O5 approaches the composites at high temperatures (T > 1000 K), it is clear the composites achieve a significantly lower thermal conductivity at room temperature, due to the scattering of long-wavelength phonons by the W-rich precipitates that span length scales from nanometers to microns. These low W-content composites have superior thermoelectric performance in every regard compared to Nb2O5; Nb2O5 is too insulating to measure electrical resistivity, whereas (W1.083O3)x(Nb2.1O5)1−x materials exhibit metallic electrical transport. Yet despite the much larger expected contribution from electrical carriers to thermal conductivity in these materials, their thermal conductivity is 30% less than that of insulating Nb2O5 at 300 K.

At high temperature in addition to the secondary phase inclusions, intrinsic defects in the crystal structure may also play a role (e.g., crystallographic shear planes, site defects from oxygen deficiency, and mass contrast from W/Nb site substitution), scattering short-wavelength phonons and reducing their ability to carry heat. Although the thermal conductivity of Nb2O5 approaches the composites at high temperatures (T > 1000 K), as it shares many of these intrinsic defects present in (W1.083O3)x(Nb2.1O5)1−x, the composites remain the materials with minimum thermal conductivity.

Taken together with electrical property measurements, thermal conductivity measurements reveal the largest zT in this compositional series is achieved by the x = 0.08 member, where zT = 0.1 at 950 K, the highest temperature measured here (Figure 10).

FIG. 10.

Physical property measurements reveal that the highest power factor and zT achieved in these samples is for the x = 0.08 member, which achieves zT = 0.1 at 950 K, the highest temperature measured here. κzT is shown to provide a normalized zT without consideration of thermal conductivity.

FIG. 10.

Physical property measurements reveal that the highest power factor and zT achieved in these samples is for the x = 0.08 member, which achieves zT = 0.1 at 950 K, the highest temperature measured here. κzT is shown to provide a normalized zT without consideration of thermal conductivity.

Close modal

Early transition metal oxides were targeted using a datamining approach, and dense pellets of (W1.083O3)x(Nb2.1O5)1−x composites were prepared by SPS, which led to simultaneous reaction and densification of WO3, Nb2O5, W, and Nb powders in 6 min at 1473 K. The samples were doped to nominally have 0.5 electrons per metal cation. Backscattered electron micrographs revealed samples with low W content (x ≤ 0.13) have W metal secondary phase inclusions with dimensions spanning from <20 nm to ∼2 μm. The average size and number of these precipitates in (W1.083O3)x(Nb2.1O5)1−x composites increases controllably with W content up to at least x = 0.13, and the dimensions of these precipitates are ideal for scattering phonons with a broad wavelength distribution. Furthermore, these inclusions are stable even after repeated thermal cycling, as monitored by the electrical properties with cycling up to the highest temperatures measured, ie. 950 K. This leads to composite materials with reduced thermal conductivity across all temperatures investigated (300 K to 1050 K), even with W content as low as 0.3 mol% (x = 0.02 in (W1.083O3)x(Nb2.1O5)1−x).

These (W1.083O3)x(Nb2.1O5)1−x composites are a promising platform for n-type oxide thermoelectric materials. Even without electronic optimization (e.g. doping studies), these composites are in the same performance class as CaMnO3-based materials, which have been heavily researched (Figure 11). Like other n-type oxides the Seebeck coefficient has a linear temperature dependence, electrical resistivity near 10−2 Ωcm, and ∼2Wm−1 K−1 at high temperatures. Physical property measurements reveal that the highest power factor and zT achieved in these samples is for the x = 0.08 member, which achieves zT = 0.1 at 950 K. Although this is the highest temperature measured here, the materials are stable to higher temperatures, where the zT should be higher. Further studies should lead to improvement in electrical transport properties, and improved thermoelectric performance in this new class of n-type oxides that are stable at high temperature.

FIG. 11.

W–Nb–O materials contextualized within the larger survey plot of current oxide thermoelectric materials reveals that W–Nb–O materials at 1000 K (X) are an exciting platform for potential n-type oxides. The (W1.083O3)x(Nb2.1O5)1−x composites studied here have comparable performance to CaMnO3 based materials, despite not having been optimized at this stage. The size of the circle is proportional to the material power factor. Dashed lines of constant power factor give bounds for reasonable thermoelectric performance, and this plot reveals the lack of high-performance n-type oxide materials. Data from.1 

FIG. 11.

W–Nb–O materials contextualized within the larger survey plot of current oxide thermoelectric materials reveals that W–Nb–O materials at 1000 K (X) are an exciting platform for potential n-type oxides. The (W1.083O3)x(Nb2.1O5)1−x composites studied here have comparable performance to CaMnO3 based materials, despite not having been optimized at this stage. The size of the circle is proportional to the material power factor. Dashed lines of constant power factor give bounds for reasonable thermoelectric performance, and this plot reveals the lack of high-performance n-type oxide materials. Data from.1 

Close modal

We thank the National Science Foundation for support of this research through NSF-DMR 1121053. MWG thanks the Natural Sciences and Engineering Council of Canada for an NSERC Postgraduate Scholarship and the US Department of State for an International Fulbright Science & Technology Award. JED is supported by the National Science Foundation Graduate Research Fellowship, under grant no. 1144085. The research carried out here made use of the shared experimental facilities of the Materials Research Laboratory: an NSF MRSEC, supported by NSF-DMR 1121053. The MRL is a member of the the NSF-supported Materials Research Facilities Network (www.mrfn.org). This work benefited from the use of the Advanced Photon Source at Argonne National Laboratory supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences (DE-AC02-06CH11357).

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