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High Seebeck Coefficient and Unusually Low Thermal Conductivity Near Ambient Temperatures in Layered Compound Yb2–xEuxCdSb2

  • Joya A. Cooley
    Joya A. Cooley
    Department of Chemistry, University of California, One Shields Avenue, Davis, California 95616, United States
  • Phichit Promkhan
    Phichit Promkhan
    Department of Chemistry, University of California, One Shields Avenue, Davis, California 95616, United States
  • Shruba Gangopadhyay
    Shruba Gangopadhyay
    Department of Chemistry, University of California, One Shields Avenue, Davis, California 95616, United States
  • Davide Donadio
    Davide Donadio
    Department of Chemistry, University of California, One Shields Avenue, Davis, California 95616, United States
    IKERBASQUE, Basque Foundation for Science, E-48011 Bilbao, Spain
  • Warren E. Pickett
    Warren E. Pickett
    Department of Physics, University of California, One Shields Avenue, Davis, California 95616, United States
  • Brenden R. Ortiz
    Brenden R. Ortiz
    Department of Physics, Colorado School of Mines, Golden, Colorado 80401, United States
  • Eric S. Toberer
    Eric S. Toberer
    Department of Physics, Colorado School of Mines, Golden, Colorado 80401, United States
  • , and 
  • Susan M. Kauzlarich*
    Susan M. Kauzlarich
    Department of Chemistry, University of California, One Shields Avenue, Davis, California 95616, United States
Cite this: Chem. Mater. 2018, 30, 2, 484–493
Publication Date (Web):December 18, 2017
https://doi.org/10.1021/acs.chemmater.7b04517

Copyright © 2017 American Chemical Society. This publication is licensed under these Terms of Use.

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Abstract

Zintl phases are promising thermoelectric materials because they are composed of both ionic and covalent bonding, which can be independently tuned. An efficient thermoelectric material would have regions of the structure composed of a high-mobility compound semiconductor that provides the “electron–crystal” electronic structure, interwoven (on the atomic scale) with a phonon transport inhibiting structure to act as the “phonon–glass”. The phonon–glass region would benefit from disorder and therefore would be ideal to house dopants without disrupting the electron–crystal region. The solid solution of the Zintl phase, Yb2–xEuxCdSb2, presents such an optimal structure, and here we characterize its thermoelectric properties above room temperature. Thermoelectric property measurements from 348 to 523 K show high Seebeck values (maximum of ∼269 μV/K at 523 K) with exceptionally low thermal conductivity (minimum ∼0.26 W/m K at 473 K) measured via laser flash analysis. Speed of sound data provide additional support for the low thermal conductivity. Density functional theory (DFT) was employed to determine the electronic structure and transport properties of Yb2CdSb2 and YbEuCdSb2. Lanthanide compounds display an f-band well below (∼2 eV) the gap. This energy separation implies that f-orbitals are a silent player in thermoelectric properties; however, we find that some hybridization extends to the bottom of the gap and somewhat renormalizes hole carrier properties. Changes in the carrier concentration related to the introduction of Eu lead to higher resistivity. A zT of ∼0.67 at 523 K is demonstrated for Yb1.6Eu0.4CdSb2 due to its high Seebeck, moderate electrical resistivity, and very low thermal conductivity.

Introduction

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Thermoelectric (TE) generators can transform thermal energy into electrical energy. This thermal energy can come from a wide variety of sources such as vehicle exhaust systems (1) and solar radiation. (2) These devices work using the Seebeck effect in two dissimilar materials. Ideally, an n-type (electron conduction) semiconductor and p-type (hole conduction) semiconductor are joined electrically in series and thermally in parallel across a temperature gradient. In a semiconductor, carriers diffuse from the hot side of the gradient to the cold side and, since the carriers are charged, allow for a voltage difference to develop. When a load is connected across the two materials, electricity can flow. The maximum efficiency is governed in part by a dimensionless figure of merit, zT (= S2T/ρκ). It follows that, in order to achieve high efficiency, it is necessary to have a high Seebeck coefficient (S), low electrical resistivity (ρ), and low thermal conductivity (κ). Slack (3) suggested that the ideal thermoelectric material would be a “phonon–glass, electron–crystal” (PGEC), in which the material conducts heat poorly, as in an amorphous glass, but can conduct electronic charge very well, as in a crystalline material.
In general, semiconductors provide a suitable combination of high thermopower and reasonably low electrical resistivity to make suitable PGEC materials, if their structures can be suitably tuned to yield low thermal conductivity. Zintl phases, (4) a subset of intermetallic phases combining ionic and covalent bonding, have been targeted as viable TE materials as they often intrinsically exhibit PGEC properties since they readily form valence precise semiconductors. Zintl phases have been studied at temperatures far above room temperature, up to ∼1250 K, for TE power generation applications, and show great promise. (5−13) Yb14MnSb11 (5) and variants thereof (14) illustrate well the PGEC concept, with a lattice thermal conductivity reaching as low as 0.16 W/m K combined with an electrical resistivity of less than 1.0 mΩ cm in Yb14MgBi11. (15) PGEC behavior has also been seen in other Zintl phases such as coinage-metal-stuffed Eu9Cd4Sb9, (9) a derivative of Mg3+δSb2, Mg3.2Sb1.5Bi0.49Te0.01, (16) and clathrate Ba8Ga16Ge30. (17)
Despite this progress, there is still considerable room for growth with respect to applications at lower temperatures, where there are a significant number of technologies that would benefit from the additional energy recovery from heat, for example, solar thermoelectric generators (STEGs). STEGs are an emerging technology that, much like photovoltaics, harvest energy from the sun to create electricity, but use concentrated thermal energy the sun provides rather than radiation. (18) With a zT = 1 and hot and cold sides of temperatures 220 and 20 °C, respectively, STEGs can reach 8.6% efficiency. (2) Thus, high zTs in this temperature range are of great consequence, and more study to find thermoelectrically competitive materials in this temperature range is worthwhile.
Discovery of materials with high zT involves some intuition (19) and appropriate choice of structure. (20) Semiconductors with large Seebeck coefficients at high temperature can be difficult to identify, thereby focusing research on low thermal conductivity materials. One method of enhancing zT involves striving for very low inherent κ in a material and working to further lower the lattice thermal conductivity (κlatt). So-called complex structures have been linked to having low κlatt because their large unit cells, use of heavy elements, and complex bonding schemes allow for significant phonon scattering, thus acting as a “phonon glass”. Some of these complex structures include clathrates and skutterudites, (21) as well as Zintl phases such as Yb14MnSb11, (5) coinage-metal-stuffed Eu9Cd4Sb9, (9) Ca5Al2Sb6, (11) and nanostructured materials. (22,23) Despite this apparent requirement of complex structure types, phases with relatively simple layered structures have demonstrated favorable thermoelectric properties: often the layered structure plays a role in the compound’s PGEC properties. For example, NaCo2O4 contains CoO2 layers, responsible for electronic conduction, while the disordered Na layer acts as a phonon scattering mechanism, allowing for κ of ∼0.5 W/m K at relatively low temperatures. (24) Bi2Te3, a state-of-the-art TE material that is in commercial use today, is a layered compound whose low thermal conductivity of near ∼0.7 W/m K resembles that of a glass and contributes to its high figure of merit near room temperature. (25) Layered Zintl phases have also garnered attention: the AM2Pn2 (A = Ca, Sr, Eu Yb; M = Zn, Cd; Pn = As, Sb) family of compounds, in the CaAl2Si2 structure type, has yielded a wealth of structures with low κ and high S, often resulting in impressive zT values. (26,27)
Herein, we have investigated the layered solid-solution Yb2–xEuxCdSb2 (Figure 1) as a potential thermoelectric material for near room temperature applications as motivated by the promising S and ρ reported for the single crystals. (28,29) Yb2–xEuxCdSb2 crystallizes in the polar space group of Cmc21. The structure is almost like that of Ca2CdSb2 except for a small distortion that destroys the center of inversion. The lack of an inversion center, which splits degeneracies compared to the higher-symmetry structure, should translate into a more complex phonon spectrum and phonon–phonon interaction and thereby lower the thermal conductivity and increase the figure of merit. Compared to Ca- or Ba-based compounds the noncentrosymmetric space group is unique to this combination of Yb and Eu. (28) The layered structure and zigzag bonding of the anionic substructure are reminiscent of SnSe, (30,31) a phase exhibiting significant anharmonicity yielding low thermal conductivity. Therefore, one might expect anharmonicity to play a role in this structure type as well. There are two crystallographic sites in this structure for the rare earth cations, providing an additional degree of freedom. Not only does site substitution provide an alloy scattering mechanism to reduce thermal conductivity, but site preferences for the two cations provide additional complexity that can tune properties. In this solid-solution phase, we are substituting Eu for Yb, and Eu preferentially substitutes on the interlayer (Yb1) site, shown in Figure 1. Eu will add additional f-bands to the band structure and, more importantly, may provide Eu–Sb bonding, as is the case in the coinage-metal-stuffed Eu9Cd4Sb9 complexes, (9) to also affect the gap or placement of the Fermi level. The structural features of Yb2CdSb2 combined with the high thermopower reported at room temperature (28) provide incentive to investigate Yb2–xEuxCdSb2 in order to confirm a high figure of merit phase based on the chemical intuition outlined above.

Figure 1

Figure 1. (a) Ball-and-stick depiction of the structures of Yb2–xEuxCdSb2 with one unit cell outlined. Yb, Eu, Cd, and Sb are shown as teal, pink, purple, and tan spheres, respectively. Coordination spheres of sites: (b) Yb1, interlayer, and (c) Yb2, intralayer.

Experimental Section

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Synthesis

Bulk samples were synthesized through a combination of mechanical milling to pulverize elements to a homogeneous powder and spark plasma sintering (SPS) to simultaneously provide a phase pure sample and densify the powder into a pellet. Yb was filed to pieces using a metal rasp and combined stoichiometrically with cut Eu pieces, Cd shot, and Sb shot (sources: Eu metal, Stanford Materials, 99.99%; Yb pieces, Metall Rare Earth Limited, 99.99%; Cd shot, Alfa Aesar, 99.95%; and Sb shot, Alfa Aesar 99.9%). The elements were weighed according to the following stoichiometry: 2 – x Yb:x Eu:1 Cd:2 Sb (x = 0.3, 0.8, 1, and 1.1) to total reaction masses of 3–3.5 g. For each reaction, Yb, Cd, and Sb, and two tungsten carbide (WC) balls (diameter 8.0 mm) were loaded into a 5 cm3 stainless steel ball mill, capped via a Viton O-ring with WC disks held in place by Teflon caps. The milling container was heat-sealed in two poly bags to prevent oxidation in air and placed in an SPEX 8000 M for 1 h total in two 30 min milling duration increments. Between increments, the milling container was transferred to a drybox, and the inside and end-caps were scraped down with a metal spatula; Eu was added for the second ball-milling. The resultant powder was transferred to a drybox and ground via alumina mortar and pestle to remove any large agglomerates and sieved through a 100 mesh sieve. Approximately 2.5–3 g of powder was transferred to a 12.7 mm diameter high-density graphite die (from POCO) for densification via a Dr. Sinter SPS-2050 spark plasma sintering (SPS) system (Sumitomo, Tokyo, Japan). The die was transferred to the instrument chamber, which was evacuated (below ∼20 Pa) and filled with a partial pressure of Ar. The temperature was increased from room temperature to 500 °C in 15 min and then to 600 °C in 1 min (to prevent temperature overshoot), and it remained stable for 45 min. When the temperature reached ∼480 °C and compression began, the force was increased from 2 to 8 kN. Afterward, samples were cooled to room temperature with pressure released. Sample densities were larger than 90% of the theoretical density according to the Archimedes method.

Powder X-ray Diffraction

Ground powders of pelleted samples were loaded on a Bruker zero-background holder for powder X-ray diffraction (XRD). XRD data were collected by a Bruker D8 Advance diffractometer with Cu Kα radiation (λ = 1.54060 Å) with 2θ range 20–80° with a step size ∼0.02° and scan rate of 1 s per step. Refinements of unit cell parameters were performed with the program FullProf. (32)

Electron Microprobe Analysis

Pellet pieces from each sample were examined for phase composition by electron microprobe analysis (EMPA) using a Cameca SX-100 electron probe microanalyzer equipped with a wavelength-dispersive spectrometer with 15 kV accelerating potential and a 20 nA beam current. Small pellet pieces (∼1.5 mm × 1.5 mm × 1.5 mm) were coated in epoxy and polished for analysis. The polished samples were mounted on 25 mm metal rounds with carbon tape and coated with carbon for conductivity. Characteristic X-rays generated by samples were qualitatively analyzed by wavelength-dispersive spectroscopy to determine compositions of samples using X-ray intensities of Yb, Eu, Cd, Sb as compared to calibrated standards Yb14MnSb11, EuGa4, Cd (metal), Sb (metal). Elemental mapping was performed to qualify element homogeneity.

Thermoelectric Property Measurement

High-temperature thermal diffusivity (λ, shown in Table S3 of the Supporting Information) was measured between 348 and 523 K on 12.7 mm diameter circular pellets using a Netzsch Laser Flash Analysis (LFA) 457 instrument. Measured thermal diffusivity was used to calculate thermal conductivity using the equation κtot = λρCp, where ρ is the density and Cp is the heat capacity. The density was measured using the Archimedes method, and the heat capacity was estimated using the Dulong–Petit equation Cp = 3RN/FW (R = 8.3145 J/mol K, N = number of moles of atoms in the formula unit, FW = formula weight). Experimental heat capacity (Figure S3 in the Supporting Information) was measured with a Netzsch Thermal Analysis DSC 404 F3 with a TASC 414/4 controller using alumina pans placed in platinum pans and the platinum furnace from 313 to 523 K. The Netzsch software package automatically subtracted the data and referenced against the standard (sapphire) to calculate the heat capacity values. A Linseis LSR-3 instrument was used to measure Seebeck coefficient and electrical resistivity via a standard four-probe method from 348 to 523 K under a helium atmosphere each sample. Samples were cut into bars (∼12 mm × 2 mm × 2 mm) using a Buehler diamond saw and polished before measurement. The probe distance was 6 mm. The Seebeck coefficient, electrical resistivity, and thermal conductivity data were fit to fifth- or sixth-order polynomial functions to calculate zT values.

Computational Details

The present DFT-based electronic structure calculations were performed using the full-potential augmented plane-wave plus local orbital method as implemented in the WIEN2K code. (33−35) The generalized gradient approximation (GGA) as described by Perdew, Burke, and Ernzerhof (PBE) was used for the underlying semilocal exchange-correlation potential. (36) GGA+U calculations with fully localized limit are used for approximating local repulsion effects on localized atomic orbitals. In this method an on-site Hubbard repulsion U is included for the rare earth 4f-orbitals and Cd 4d-orbitals. The double counting correction approach proposed by Anisimov et al. (37) and Liechtenstein et al. (38) was chosen for these highly localized orbitals, and the value U = 8 eV and Hund’s exchange parameter J = 1 eV were chosen for f-orbitals in Yb and Eu. For Cd d-orbital we applied U = 3 eV. Although the Cd 4d-orbital shows minor participation near the valence edge, Flage-Larsen et al. (39) have shown that, in the case of YbZn2Sb2, applying an on-site U for Zn d states can lead to a more accurate description of electronic structure; therefore, the on-site U for Cd was employed. Also, to keep numerical consistency with magnetic YbEuCdSb2 we applied spin polarized calculations for both Yb2CdSb2, nonmagnetic (4f14 configuration), and YbEuCdSb2, magnetic (Eu2+, 4f7 configuration). The wave functions in the interstitial regions were expanded in plane-waves up to a cutoff of RmtKmax = 9 (Rmt is the smallest muffin tin radius of atoms, and Kmax is the maximum plane-wave vector) to achieve convergence for energy eigenvalues (especially the 4f-bands). Nonoverlapping atomic sphere radii, in a.u., were chosen as 2.5 for all the constituent atoms. Plane-waves with a kinetic energy cutoff of 500 eV are used as basis sets, and a 4 × 1 × 2 k-point grid in Brillouin zone is chosen as per the Monkhorst–Pack scheme. All the pseudopotentials were calibrated with respect to the density of states obtained from all-electron WIEN2K calculations, and for both cases the same Hubbard U was applied for all rare earth 4f-orbitals. Applied pseudopotentials have the following valence configurations: 5s2, 5p6, 4f14, 6s2 for Yb; 5s2, 5p6, 4f7, 6s2 for Eu; 4d10, 5s2 for Cd; and 5s2, 5p3 for Sb. The lattice structure relaxations are carefully performed so that the forces on atoms are smaller than 0.2 meV/Å, in which the conjugate gradient algorithm is utilized.
Starting from the experimental crystallographic coordinates, (29) the crystal lattice parameters were optimized, computing the energy versus volume curves of both compounds and fitting the Vinet equations of state. (40) The electronic transport properties (e.g., Seebeck coefficient and electrical conductivity) were calculated by solving the semiclassical Boltzmann transport equation assuming constant relaxation time approximation (CRTA) as implemented in the BoltzTrap package, (41) using band energies obtained from WIEN2K electronic structure calculation as input. A dense mesh of 31 × 8 × 16 k-points in the irreducible wedge of the Brillouin zone was used to obtain converged results.
Vibrational analysis and phonon dispersion curves were computed using the plane-wave pseudopotential approach, as implemented in the VASP package, (42) with projector-augmented-wave pseudopotentials. (43) GGA+U applied for frozen phonon calculations is based on the approach by Dudarev et al. as implemented in VASP. (44) Phonon dispersion curves, speed of sounds calculations, and phonon density of states are obtained using the supercell frozen phonon method and the phonopy code. (45)

Speed of Sound Measurements

Speed of sound measurements are performed using an Olympus 5072PR Pulser/Receiver system with a gain of 20 dB and a 5 kHz signal. Both longitudinal and shear measurements were made, using Olympus V112 (longitudinal) and Olympus V156 (shear) transducers and an Atten ADS 1102 oscilloscope.

Hall Measurements

Hall effect and resistivity measurements are performed using the Van der Pauw geometry on a home-built apparatus. (46) Measurements are conducted up to 523 K under dynamic vacuum (<10–4 Torr). Contacts are pressure-assisted nichrome wire.

RESULTS/DISCUSSION

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Powder X-ray Diffraction and Sample Purity

The Yb2CdSb2 structure type crystallizes in the noncentrosymmetric space group Cmc21 and has been well-described by previous publications. (28,29) Ball-milling of elements followed by spark plasma sintering (SPS) resulted in samples with densities greater than 90% of theoretical values. Figure 2 shows the powder X-ray diffraction (PXRD) patterns of ground pellets of Yb2–xEuxCdSb2 (x = 0, 0.36, 0.85, 1.13, and 1.18, determined from EMPA provided below) compared to the calculated pattern for Yb2CdSb2 (top). Black asterisks indicate small impurity peaks arising from impurity phases, one of which disappears as Eu content increases (∼25°) that could not be indexed to this structure type. This peak does match to a peak from the calculated pattern of YbCd2Sb2, which has also been found in a highly defected form by electron microprobe analysis (Yb1.06(4)Cd1.93(2)Sb2.00(3)); however, there are too few peaks visible to unequivocally assign this phase as the impurity by PXRD. Since one impurity (∼29°) is a consistent impurity for the series, we expect this study will provide insight into the intrinsic properties of the system. (47) Samples show a gradual left-shift in peak position with increasing x as unit cell parameters increase, consistent with Eu being incorporated into the structure primarily in the Yb1 six-coordinate site (Figure 1b) as Eu2+ has a larger ionic radius than Yb2+ (1.14 and 1.03 Å, respectively (48)). Single crystal X-ray diffraction refinement results from Yb2–xEuxCdSb2 indicate that Eu2+ prefers the Yb1 site, consistent with these results. (28,29) The Yb1 site, with its greater average bond distance of 3.338 Å, accommodates the larger Eu2+ more easily than the Yb2 site with its average bond distance of 3.168 Å. Lattice parameters extracted from whole-pattern fitting show increasing unit cell volume with increasing Eu concentration as expected, in good agreement with lattice parameters reported from single crystal X-ray diffraction results (Supporting Information). (28,29) No further Rietveld analysis was employed due to significant peak broadening, likely from strain stemming from the ball-milling process as well as the high force applied by the SPS process.

Figure 2

Figure 2. Powder X-ray diffraction (PXRD) patterns for synthesized samples, Yb2–xEuxCdSb2. Black asterisks indicate an impurity unable to be indexed by PXRD. Whole-pattern fitting for x = 0, 0.36, 0.85, 1.13, and 1.18 revealed unit cell volumes of 581, 586, 600, 607, and 610 Å3, respectively.

EMPA was used to verify stoichiometry and qualify the element homogeneity (EMPA images can be found in the Supporting Information). The average compositions of samples as determined by wavelength-dispersive spectroscopy (WDS) are shown in Table 1. These results indicate that the Yb2–xEuxCdSb2 phase was formed in each case. Results are also in relatively good agreement with the preparative amounts of Eu and Yb and show, within error, that the total cation count is near 2 in all samples. The samples of x = 1, 1.1 show cation totals with large standard deviations which can be attributed to inhomogeneity due to grain boundaries. The experimental value of x is employed to identify samples throughout.
Table 1. Sample Compositions of Yb2–xEuxCdSb2 As Determined by Electron Microprobe Analysis (EMPA)
preparative Eu amountexptl Eu amountsample composition (as determined by EMPA)total cation count (Eu + Yb)
00Yb1.99(3)Cd1.02(3)Sb1.99(1)1.99(3)
0.30.36Yb1.58(2)Eu0.36(1)Cd1.03(3)Sb2.017(7)1.94(2)
0.80.85Yb1.18(6)Eu0.85(3)Cd0.98(1)Sb1.98(3)2.03(3)
11.13Yb0.98(7)Eu1.13(7)Cd0.89(11)Sb1.99(4)2.11(10)
1.11.18Yb0.87(4)Eu1.18(3)Cd0.93(6)Sb2.01(1)2.05(6)

Electronic Structure

Yb2CdSb2 is a valence precise compound that can be described by the Zintl formalism. In a simple Zintl picture, Yb2+ donates two electrons to the more electronegative atoms, as confirmed by magnetic measurements. (28) The [CdSb2]4– layers can be considered as composed of Cd2+ and Sb3–; thus, [Yb2+]2[Cd2+][Sb3–]2 is considered to be a charge balanced, (49) narrow-gap semiconductor. Given the inclusion of d- and f-block metals, the definition of Zintl phases has expanded, and in this case, Cd acts as a metalloid bonding with Sb. Density functional theory (DFT) calculations were employed to provide insight into the electronic structure, electronic transport and vibrational properties of these materials.

a. Crystal Structure

Electronic structure calculations were performed for two stoichiometric configurations of Yb2–xEuxCdSb2 with x = 0 and 1. Consistent with experimental structural analysis, the calculations support the preference of Eu substitution in the Yb1 site. In fact, the electronic structure would not converge when Eu substitutes on the Yb2 sites, without Eu occupying the Yb1 sites. Conversely, the system for x = 1.25 (that is, all Yb1 sites plus one Yb2 site out of four are substituted with Eu) converges and is stable. Theoretical lattice parameters are in good agreement with experiments (see Table S1 in the Supporting Information). In particular, the calculated volume of Yb2CdSb2 (586 Å3) is within 1.6% of the experimental estimate (577 Å3). For x = 1 Eu substitution induces a significant volume expansion due to the larger radius of Eu2+ (calculated 620 Å3; experimental 607 Å3). Further substantial volume expansion occurs for x = 1.25 (641 Å3). The lack of such drastic volume expansion in experiments indicates that for x > 1 the Eu atoms are not being incorporated on the Yb sites as a solid solution. The additional Eu may simply be present at the grain boundary. Electronic structure and electronic transport property calculations were performed using fully optimized crystal structures.

b. Electronic Structure

Spin polarized density functional theory calculations show that Yb2CdSb2 is nonmagnetic, as expected due to the presence of closed shell Yb2+ (4f14), whereas the magnetic moment of YbEuCdSb2 stems from the Eu 4f-orbitals. (28) Eu2+ in the crystal retains its atomic magnetic moment (6.94 μB in the atomic sphere) of 7 unpaired electrons, suggesting that 4f-orbitals are inert and participate neither in chemical bonding nor in the Fermi level states involved in charge transport. We show below that this picture is oversimplified.
The band structure of Yb2CdSb2 is shown in Figure 3a, and both spin-up (—) and spin-down (---) bands of Yb2–xEuxCdSb2 (x = 1) are shown in Figure 3b. The valence band maximum occurs at the Γ point while the conduction band minimum occurs at the Y symmetry point, yielding an indirect band gap (92 meV) in Yb2CdSb2. This finding of a gap differs from the report of Xia and Bobev [ref (29)], who used the augmented spherical wave method that applies spherical averages to potential and density within the atom sphere. Also not applied in ref (29) were the on-site repulsive interactions on Yb. The distinction results in our obtaining a band gap, and in placing the Yb 4f levels at −1.9 eV in our calculaion, whereas their result had the center of the Yb 4f-bands centered at −0.2 eV, leading to strong hybridization with the conduction bands up to the hole chemical potential at the higher hole concentrations.

Figure 3

Figure 3. Band structure of Yb2–xEuxCdSb2 with valence (---) and conduction (—) bands indicated (a) for x = 0, and (b) for combined spin majority (—) and spin minority (---) bands for x = 1. Panel a shows flat bands at −1.9 eV from Yb 4f states whereas panel b shows flat bands around −1.9 eV arising from Yb 4f complemented with Eu 4f-bands at −2.6 eV. Panel c presents an enlargement of the x = 0 band in panel a near the band gap.

The calculated band gap opens slightly to 111 meV for YbEuCdSb2. While band gaps calculated with DFT may be underestimated, the opening of the gap with increasing Eu concentration is consistent with the literature and experimental results presented below. Band gaps calculated from maximum Seebeck coefficient (Eg = 2eSmaxTmax) in YbCd2Sb2 and EuCd2Sb2 (50) and with band structure calculations for YbMg2Bi2 and EuMg2Bi2 provide a similar trend. (51) In both compounds the flat lanthanide f-bands within the band structure can be clearly identified. In Yb2CdSb2 the valence band edge contains one roughly isotropic band, with a second band lying around 150 meV higher binding energy. An enlarged view is provided in Figure 3c. The valence band edge is very similar for both compounds, but the spin-down channel in Figure 3b reveals an exchange splitting of 0.5 eV reflecting the exchange coupling between Eu 4f- and 5d-orbitals. The half filled 4f-orbitals provide an additional magnetic scattering mechanism when Eu is substituted into Yb2CdSb2.
The projected density of states (PDOS, Figure 4) for both x = 0 and x = 1 shows that the main character of the valence band is dictated by Sb 5p states, with Sb p, Cd p, and Yb f-orbitals contributing at higher hole binding energy. In both compounds, Sb 5p, Cd 5p, lanthanide 5d electrons are the majority participants in covalent bonding in the valence band (see more details in Figure S8 in the Supporting Information). While the band structure (Figure 3) might suggest little interaction of the f-orbitals, the PDOS reveals Yb f character rising rapidly a few tenths of an eV below the gap. In the x = 1 compound the f-bands of Eu, which lie −2.6 eV below the gap, push the Yb f character near the gap to higher energy, thus accentuating the f character of holes doped into the valence band. Such doping-induced changes in the PDOS are small enough not to produce major changes in hole mobility, but dictate the different trends in the Seebeck coefficient computed for pure and Eu-doped Yb2CdSb2.

Figure 4

Figure 4. Orbital-projected density of states plot using the DFT+U functional for Yb2–xEuxCdSb2. (a) Participating atomic orbitals near the valence edge are shown for x = 0. (b) All participating orbitals near the valence edge are shown for x = 1. The majority spin channel is shown for x = 1 in these plots. Additional detailed PDOS plots are provided in the Supporting Information.

Electronic Transport Properties

On the basis of Zintl valence-counting rules and consistent with the DFT-calculated band structure calculations, stoichiometric Yb2–-xEuxCdSb2 compounds are expected to show semiconducting behavior. The electrical resistivities (Figure 5a) of x = 0, 0.36 and x = 0.85 samples are small and show metallic temperature dependences, as in degenerate semiconductors, while x = 1.13 and x = 1.18 show higher resistivity, consistent with both the widening of the band gap and the lowered carrier concentrations. The transition from degenerate to nondegenerate behavior seen in electrical resistivity is reflected in Seebeck measurements, as it would be expected that metallic systems have lower Seebeck coefficients. S(T, x) increases monotonically with Eu concentration, correlating with the system becoming less metallic. Hole concentrations at 323 K (Table 2) are in the range of ∼1019 due to intrinsic and compositional defects. Increasing Eu concentration causes the hole concentration to decrease, possibly because there are more defects on Yb sites, as has been seen in other cases of Eu substitution for Yb. (52,53) This is consistent with the cation totals seen in the composition (Table 1). According to the Nagle electronegativity scale, (54) Yb is more electronegative than Eu which has been suggested to influence carrier concentration because of incomplete electron donation from the cation. (53) The 373 K mobility data (Table 2) show high values for x = 0, 0.36 and 0.85, not often seen in Zintl phases, and values sharply decrease for x = 1.13 and x = 1.18. Additional Hall data are provided in the Supporting Information.

Figure 5

Figure 5. Temperature dependent (a) electrical resistivity and (b) Seebeck coefficient on polycrystalline samples of Yb2–xEuxCdSb2. The calculated Seebeck coefficients (open symbols) use experimental carrier (hole) concentration from Table 2. For x = 1 the carrier concentration is obtained by interpolating between the experimental carrier concentrations at x = 0.85 and x = 1.13. Experimental data points are shown as filled symbols and fits as solid lines.

Table 2. Carrier (Hole) Concentration and Mobility Data for Each Sample Composition, Yb2–xEuxCdSb2, at 323 K
Eu conc (x)carrier conc (n) (cm–3)mobility (cm2/(V s))
04.53 × 101970.30
0.363.56 × 101972.18
0.858.96 × 101895.93
1.135.37 × 101824.55
1.181.41 × 101917.51
The x = 0.36 phase shows the highest carrier concentration. From this, we can propose that Yb vacancies are the major source for charge carriers (total cation stoichiometry is less than 2, Table 1), and substituting the Yb site with other alio- and isovalent cations can be a useful tool for band engineered thermoelectrics. The high mobility for the samples with x < 1 makes these phases attractive for further optimization. The nonmonotonous mobility is one of the key subjects of interest of further investigation.
Experimental temperature dependent Seebeck coefficients are shown in Figure 5b (filled symbols). All are positive in the entire range, indicative of holes as the dominant charge carriers. Consistent with previous studies on single crystal samples, the Seebeck is large in all samples. The Seebeck coefficient increases with Eu concentration, reaching a maximum value of ∼270 μV/K for the x = 1.18 phase at 523 K. Each sample shows nearly linear temperature dependence, typical of heavily doped semiconductors. Additionally, theoretically computed Seebeck coefficients are plotted in Figure 5b (open symbols) for x = 0, 0.85, and 1. To compare to experiments, which are performed on polycrystalline samples, S is defined as 1/3 of the trace of the anisotropic Seebeck tensor of a single crystal. For these calculations, the experimentally determined carrier concentrations (Table 2) were used, and for the x = 1 compound the carrier concentration was interpolated from the experimental data at x = 0.85 and 1.13. To compute the Seebeck coefficient of Yb1.15Eu0.85CdSb2 the band structure of YbEuCdSb2 was employed. Theory and experiments for this system exhibit excellent agreement, which suggests all electron-based DFT+U level of calculations and the semiclassical Boltzmann transport equation in CRTA can provide a faithful prediction of Seebeck coefficients for the Yb2–xEuxCdSb2 (2-1-2) Zintl family.
The effect of carrier concentration (n) on the calculated Seebeck coefficients is shown in Figure 6, which reveals that Yb2CdSb2 exhibits a crossover in its temperature dependence, at n ∼ 7 × 1019 cm–3. Below this value, the Seebeck coefficient decreases with increasing temperature, while at larger n the Seebeck coefficients become proportional to the temperature. This is the fingerprint of a transition from a nondegenerate to a degenerate semiconductor. For YbEuCdSb2S(T) is proportional to the temperature irrespective of carrier concentration. All experimental samples follow the trend of YbEuCdSb2 shown in Figure 6, so even the smallest amount of Eu substituted into the structure leads to S(T) that trends according to YbEuCdSb2 (Yb2–xEuxCdSb2, x = 1 model). From the calculation of the electrical transport coefficients of Yb2CdSb2 and YbEuCdSb2 in CRTA we can infer the ratio between the electrical resistivity (ρ) of the two systems at the same n, without making any assumption on the carrier lifetime. Since ρ is inversely proportional to the effective mass (m*) we can also estimate the change in m* upon Eu doping. We also calculated valence band effective mass m* from the second derivative of the band dispersion curve along the Γ-Y, Γ-SM, and Γ-S directions (Figure 3a–c). Since both Yb2–xEuxCdSb2 (x = 0, 1) exhibit multivalley band dispersion along three different band directions, we calculated the average (geometric mean) m* along these three directions. The results are m* = 0.33 and 0.36 for the x = 0 and 1 phases, respectively. Such a small difference in m* suggests that the changes observed in the transport coefficients as a function of Eu concentration are dictated by the carrier concentration.

Figure 6

Figure 6. Dependence of the calculated Seebeck coefficients on carriers (holes) concentration for Yb2–xEuxCdSb11 (x = 0, 1). For x = 0 configuration, a temperature dependent crossover is observed near 7 × 101 9 cm–3.

The small changes in hole mobility, which stem from the valence band edge being relatively insensitive to Eu alloying, suggest that codoping with other elements to achieve a disordered Zintl phase might be used to engineer a higher Seebeck and possibly better zT, as observed for Ca doping in (Eu0.5Yb0.5)1–xCaxMg2Bi2. (51)

Thermal Transport Properties

The total thermal conductivity, shown in Figure 7, is calculated from the thermal diffusivity (D). The DFT-calculated heat capacity of Yb2CdSb2 (Supporting Information) yields a value of 0.17 J/g K, in good agreement with values predicted by the Dulong–Petit model as well as experimentally measured Cp. The calculated thermal conductivity is extremely low and comparable to other systems with anomalously low κT such as coinage-metal-stuffed Eu9Cd4Sb9 (9) and Yb2Mn4.2Sb9. (55) The total thermal conductivity (κT) is a combination of lattice (κL) and electronic (κe) contributions. κe can be estimated using the Wiedemann–Franz relation (κe = LT/ρ). Here, L was calculated using eq 2, (56) assuming a single parabolic band (SPB) model and primarily acoustic phonon scattering. L values are shown in the Supporting Information.
(1)

Figure 7

Figure 7. Total and lattice thermal conductivity of Yb2–xEuxCdSb2. Total thermal conductivity is shown as solid symbols, and lattice thermal conductivity is shown as open symbols.

In all systems, κe is small, and subtracting it from κT yields κL (Supportng Information), the main contributor to κT. Cahill et al. (57) proposed an empirical expression for the theoretical minimum for κL as
(2)
where kB is the Boltzmann constant, V is the average volume per atom, and νT and νL are the transverse and longitudinal speeds of sound, respectively. kmin values are calculated for each composition and shown compared to κL, and in each composition, κL falls below its theoretical minimum. In a physical description, κL cannot be lower than κmin, and it can be inferred that the SPB model is inadequate or there are scattering mechanisms other than acoustic phonon scattering that are present (e.g., ionized impurity scattering) that yield unphysical values of L. Regardless, κT is extremely low at these near ambient temperatures, comparable to or lower than other materials (55,58,59) at high temperatures where κT should be at its lowest, which makes this a promising system for thermoelectric applications.
Such low κT values are not necessarily expected in such a simple structure. Similarly, PbTe, a simple rocksalt structure with no nanoinclusions, complex unit cell, or “rattler” atoms, yielded an inexplicably low κT (∼2 W/m K). This was attributed to avoided crossing in the phonon band structure and softening of certain branches generating soft lattice behavior. (60) To elucidate the possible origin of the low thermal conductivity of these compounds, the phonon band structure for x = 0 was calculated and is shown in Figure 8. The unit cell of Yb2CdSb2 contains 20 atoms, yielding 3 acoustic and 57 optical branches. The acoustic modes are confined to low frequency, below 2 THz (∼8.3 meV), with low group velocities over the whole Brillouin zone. Optical branches start as low as 0.7 THz (∼2.9 meV) and intersect with acoustic branches providing a large density of 3-phonon scattering channels. The optical bands have very low group velocities, thus contributing very little to the overall thermal conductivity. These features work in tandem to lower the lattice thermal conductivity of the compound, as has been seen in other Cd-based Zintl compounds, especially in the antimonides. (61)

Figure 8

Figure 8. Phonon dispersion of Yb2–xEuxCdSb2x = 0 by frozen phonon approximation. All optical bands are very flat in nature.

The lattice thermal conductivity can be expressed in the following equation:
(3)
where νs is the phonon group velocity, cs is the speed of sound, and lph is the phonon mean free path. According to eq 3, thermal conductivity is proportional to the speed of sound in the material. The computed longitudinal group velocity for Yb2CdSb2 along three different crystallographic axes (a, b, and c) provides speeds of sound along different axis values: 3520 m/s (a), 2644 m/s (b), 2671 m/s (c). With the average value of above three velocities, 2900 m/s as longitudinal speed of sound is obtained. The transverse (shear) phonon mode yields 1878 m/s (a), 2163 m/s (b), 2105 m/s (c), and 2048 m/s as an average. These values are in good agreement with the experimental values (Figure 9) and support the decreasing thermal conductivity with increasing Eu incorporation as this would increase point scattering throughout the system.

Figure 9

Figure 9. Shear and longitudinal speeds of sound for each sample. Speed of sound data are tabulated in Table S2 of the Supporting Information.

Additionally, as aforementioned, Eu preferentially substitutes on the interlayer site between polyanionic sheets. As the Eu concentration surpasses unity in x = 1.13 and 1.18, Eu is also found on the intralayer site (28) and stretches out the accordion-like polyanionic units. This yields disorder and strain in the lattice, further lowering thermal conductivity as can be seen by the discontinuity in κT between x = 0.85 and x = 1.13. Experimentally determined speeds of sound (Figure 9) support this conclusion as these values are comparable to other materials with low thermal conductivity, (9,55) decreasing with increasing Eu concentration. Speed of sound values suddenly make a dramatic drop at x = 1.13, further supporting the idea that the additional Eu on the intralayer site significantly affects the lattice. Thus, it can be concluded that the site specific substitution, alloy scattering, and lack of center of symmetry in this structure type lead to the low sound velocity and significant phonon–phonon scattering that are the origins of the extremely low κL.

Figure of Merit

The figure of merit of Yb2–xEuxCdSb2 (x = 0, 0.36, 0.85, 1.13, 1.18) is shown in Figure 10 as a function of temperature. All zT curves increase with temperature to 523 K with decreasing zT as Eu concentration increases. The reduction of zT is due to the large increase in electrical resistivity rendered upon having concentrations of Eu equal to or greater than that of Yb in the structure. The x = 0.36 composition of Yb2–xEuxCdSb2 shows near optimal carrier concentration and good mobility making it the best in this series for moderate to low-temperature applications, as the zT curve for x = 0.36 is slightly greater than that of TAGS, a well-known moderate temperature thermoelectric material. (62) Overall, x = 0.36 reaches a maximum zT of ∼0.7 at 523 K, an excellent temperature range for waste heat recovery such as in the case of midtemperature STEGs. However, the thermal conductivity is not the lowest for this composition, suggesting that further tuning of the 2-1-2 family can lead to increased efficiency.

Figure 10

Figure 10. Calculated figure of merit of Yb2–xEuxCdSb2.

Conclusion

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In this paper we have described a promising thermoelectric structure type, Yb2–xEuxCdSb2, with favorable thermoelectric properties. Band structure calculations reveal the compounds to be indirect band gap semiconductors with a band gap that increases with Eu concentration. Though their band centers lie well below the gap, the f-orbitals of Yb influence somewhat the values of the Seebeck coefficients through hybridization in states just below the gap, contributing to the high Seebeck coefficients observed in these compounds. Complexities in the phonon dispersion along with point defect scattering from Eu substitution allow for very low thermal conductivity. This paired with low electrical resistivity in Yb1.7Eu0.3CdSb2 results in high zT of ∼0.7 at 523 K. These results represent an initial set of measurements that guide our understanding of how to optimize this system further. The very low thermal conductivity of the series is optimized, as the lattice contribution (the major contribution to thermal conductivity) is at or below the theoretical minimum, and further efforts to improve zT can be focused on improving electrical transport properties either through controlling carrier concentration (e.g., aliovalent doping), as zT could likely be greater with higher carrier concentration, or through enhancing thermopower (e.g., band engineering). Further efforts of substituting on the rare earth sites of Yb2–xEuxCdSb2 with nonmagnetic group I or group II materials may provide a higher Seebeck coefficient and increase zT further.

Supporting Information

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The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b04517.

  • Comparison of single crystal and powder X-ray diffraction unit cell parameters, electron microprobe analysis elemental mapping images, heat capacity data, mobility and carrier concentration data, Lorenz numbers, lattice thermal conductivity with minumum values, k-path, additional projected density of states plots, Seebeck versus chemical potential, optimized DFT lattice parameters, speed of sound, thermal diffusivity, and optimized Cartesian coordinates (PDF)

Terms & Conditions

Most electronic Supporting Information files are available without a subscription to ACS Web Editions. Such files may be downloaded by article for research use (if there is a public use license linked to the relevant article, that license may permit other uses). Permission may be obtained from ACS for other uses through requests via the RightsLink permission system: http://pubs.acs.org/page/copyright/permissions.html.

Author Information

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  • Corresponding Author
    • Susan M. Kauzlarich - Department of Chemistry, University of California, One Shields Avenue, Davis, California 95616, United StatesOrcidhttp://orcid.org/0000-0002-3627-237X Email:
  • Authors
    • Joya A. Cooley - Department of Chemistry, University of California, One Shields Avenue, Davis, California 95616, United States
    • Phichit Promkhan - Department of Chemistry, University of California, One Shields Avenue, Davis, California 95616, United States
    • Shruba Gangopadhyay - Department of Chemistry, University of California, One Shields Avenue, Davis, California 95616, United States
    • Davide Donadio - Department of Chemistry, University of California, One Shields Avenue, Davis, California 95616, United StatesIKERBASQUE, Basque Foundation for Science, E-48011 Bilbao, SpainOrcidhttp://orcid.org/0000-0002-2150-4182
    • Warren E. Pickett - Department of Physics, University of California, One Shields Avenue, Davis, California 95616, United States
    • Brenden R. Ortiz - Department of Physics, Colorado School of Mines, Golden, Colorado 80401, United States
    • Eric S. Toberer - Department of Physics, Colorado School of Mines, Golden, Colorado 80401, United States
  • Notes
    The authors declare no competing financial interest.

Acknowledgments

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We thank Nicholas Botto for microprobe analysis, GAANN (J.A.C.), and NSF DMR-1405973, -1709382, and NSF CAREER award DMR-1555340 for funding. W.E.P. was supported by DOE NNSA Grant DE-NA0002908. The National Energy Research Scientific Computing Center (NERSC), a DOE Office of Science User Facility supported by the Office of Science of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231, as well as an in-house computational cluster at the University of California Davis are gratefully acknowledged.

References

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  • Abstract

    Figure 1

    Figure 1. (a) Ball-and-stick depiction of the structures of Yb2–xEuxCdSb2 with one unit cell outlined. Yb, Eu, Cd, and Sb are shown as teal, pink, purple, and tan spheres, respectively. Coordination spheres of sites: (b) Yb1, interlayer, and (c) Yb2, intralayer.

    Figure 2

    Figure 2. Powder X-ray diffraction (PXRD) patterns for synthesized samples, Yb2–xEuxCdSb2. Black asterisks indicate an impurity unable to be indexed by PXRD. Whole-pattern fitting for x = 0, 0.36, 0.85, 1.13, and 1.18 revealed unit cell volumes of 581, 586, 600, 607, and 610 Å3, respectively.

    Figure 3

    Figure 3. Band structure of Yb2–xEuxCdSb2 with valence (---) and conduction (—) bands indicated (a) for x = 0, and (b) for combined spin majority (—) and spin minority (---) bands for x = 1. Panel a shows flat bands at −1.9 eV from Yb 4f states whereas panel b shows flat bands around −1.9 eV arising from Yb 4f complemented with Eu 4f-bands at −2.6 eV. Panel c presents an enlargement of the x = 0 band in panel a near the band gap.

    Figure 4

    Figure 4. Orbital-projected density of states plot using the DFT+U functional for Yb2–xEuxCdSb2. (a) Participating atomic orbitals near the valence edge are shown for x = 0. (b) All participating orbitals near the valence edge are shown for x = 1. The majority spin channel is shown for x = 1 in these plots. Additional detailed PDOS plots are provided in the Supporting Information.

    Figure 5

    Figure 5. Temperature dependent (a) electrical resistivity and (b) Seebeck coefficient on polycrystalline samples of Yb2–xEuxCdSb2. The calculated Seebeck coefficients (open symbols) use experimental carrier (hole) concentration from Table 2. For x = 1 the carrier concentration is obtained by interpolating between the experimental carrier concentrations at x = 0.85 and x = 1.13. Experimental data points are shown as filled symbols and fits as solid lines.

    Figure 6

    Figure 6. Dependence of the calculated Seebeck coefficients on carriers (holes) concentration for Yb2–xEuxCdSb11 (x = 0, 1). For x = 0 configuration, a temperature dependent crossover is observed near 7 × 101 9 cm–3.

    Figure 7

    Figure 7. Total and lattice thermal conductivity of Yb2–xEuxCdSb2. Total thermal conductivity is shown as solid symbols, and lattice thermal conductivity is shown as open symbols.

    Figure 8

    Figure 8. Phonon dispersion of Yb2–xEuxCdSb2x = 0 by frozen phonon approximation. All optical bands are very flat in nature.

    Figure 9

    Figure 9. Shear and longitudinal speeds of sound for each sample. Speed of sound data are tabulated in Table S2 of the Supporting Information.

    Figure 10

    Figure 10. Calculated figure of merit of Yb2–xEuxCdSb2.

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      NASA MMRTG Reference Data.

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