Enhancing Thermoelectric Performances of Bismuth Antimony

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Enhancing Thermoelectric Performances of Bismuth Antimony Telluride via Synergistic Combination of MultiScale Structuring and Band Alignment by FeTe2 Incorporation Weon Ho Shin, Jong Wook Roh, Byungki Ryu, Hye Jung Chang, Hyun Sik Kim, Soonil Lee, Won-Seon Seo, and Kyunghan Ahn ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b18451 • Publication Date (Web): 05 Jan 2018 Downloaded from http://pubs.acs.org on January 5, 2018

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ACS Applied Materials & Interfaces

Enhancing Thermoelectric Performances of Bismuth Antimony Telluride via Synergistic Combination of Multi-Scale Structuring and Band Alignment by FeTe2 Incorporation

Weon Ho Shin,†,# Jong Wook Roh,‡,# Byungki Ryu,§ Hye Jung Chang,∥ Hyun Sik Kim,‡ Soonil Lee,† Won Seon Seo,† and Kyunghan Ahn*,⊥

†

Energy Materials Center, Energy & Environment Division, Korea Institute of Ceramic

Engineering & Technology, Jinju 52851, Republic of Korea. ‡

Materials R&D Center, Samsung Advanced Institute of Technology, Samsung Electronics,

Suwon 16419, Republic of Korea. §

Thermoelectric Conversion Research Center, Creative and Fundamental Research Division,

Korea Electrotechnology Research Institute, Changwon 51543, Republic of Korea. ∥

Advanced Analysis Center, Korea Institute of Science and Technology, Seoul 02792,

Republic of Korea. ⊥

Department of Chemistry, Chung-Ang University, Seoul 06974, Republic of Korea.

Keywords: thermoelectric, BST, FeTe2, band alignment, nano-precipitates, multi-scale structuring

#

Equally contributed to this work.

*To whom correspondence should be addressed. 1

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ABSTRACT It has been in difficulty in forming well-distributed nano- and meso-sized inclusions in a Bi2Te3-based matrix and thereby realizing no degradation of carrier mobility at interfaces between matrix and inclusions for high thermoelectric performances. Herein, we successfully synthesize multi-structured thermoelectric Bi0.4Sb1.6Te3 materials with Fe-rich nanoprecipitates and sub-micron sized FeTe2 inclusions by a conventional solid state reaction followed by melt-spinning and spark plasma sintering that could be a facile preparation method for scale-up production. This study presents a bismuth antimony telluride based thermoelectric material with a multi-scale structure whose lattice thermal conductivity is drastically reduced with a minimal degradation on its carrier mobility. This is possible because a carefully chosen FeTe2 incorporated in the matrix allows its interfacial valence band with the matrix to be aligned, leading to a significantly improved p-type thermoelectric power factor. Consequently, an impressively high thermoelectric figure of merit ZT of 1.52 is achieved at 396 K for p-type Bi0.4Sb1.6Te3-8 mol.% FeTe2, which is 43 % enhancement in ZT compared to the pristine Bi0.4Sb1.6Te3. This work demonstrates not only the effectiveness of multi-scale structuring for lowering lattice thermal conductivities, but also the importance of interfacial band alignment between matrix and inclusions for maintaining high carrier mobilities when designing high performance thermoelectric materials.

2


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INTRODUCTION Renewable energy sources, which are alternatives to fossil fuel as well as means of efficient energy management in existing energy sectors, have been prerequisite in order to meet ever increasing global energy demands and environmental standards at the same time. Power generation by thermoelectricity in areas where waste heat is abundant is one of the possible ways to make use of the renewable energy sources while an environmental-friendly refrigeration can be achieved by using the thermoelectric (TE) cooling.1 Thermoelectric refrigeration has several advantages over conventional vapor-compression cooling system such as a zero emission of ozone-depleting gases (no need for harmful refrigerant), a silent operation (no mechanically moving parts) and high reliability, but the main obstacle for the TE refrigeration is its low efficiency caused by inefficient TE materials.2-5 The TE efficiency of a material is determined by a dimensionless figure of merit ZT = σS2T/κ, where σ, S, T, and κ are the electric conductivity, Seebeck coefficient, temperature in Kelvin, and the thermal conductivity, respectively. It has been difficult to improve the power factor PF (= σS2) of a TE material because its σ and S vary in opposite direction through carrier concentration. On the other hand, the κ of a TE material is mainly determined by a sum of its lattice thermal conductivity (κlat) and its electronic thermal conductivity (κel). The κel is estimated according to the Wiedemann-Franz relationship κel = LσT, where L is the Lorenz number, so the κ value of a TE material is partly related to its σ value. Consequently, a significant improvement in ZT is particularly challenging because of the strong correlation between σS2 and κ. A substantial reduction in κlat by nanostructuring,6-9 all-scale hierarchical structuring,10-11 and searching for new TE materials4, 12-13 having intrinsically low κlat has been recently reported, which boosts ZT up to ~2.5 for PbTe-SrTe system11. The enhancement in PF has been also accomplished through the improvement in S by electronic band structure engineering with 3



resonance level formation,14-15 band convergence,11,

16-17

carrier filtering,18 and quantum

confinement19. Bi2Te3-based materials are known to be the best commercialized TE materials operating near room temperature.2 Various approaches have been suggested to achieve a high ZT of Bi2Te3-based materials.20-30 Renowned examples showing impressively high ZT values in a single material system are the nanostructured bismuth antimony telluride bulk alloy,7, 2021

the doped nanobulk prepared by scalable bottom-up process,24 and the liquid-phase

sintered Bi0.5Sb1.5Te3 with dense dislocation arrays embedded in grain boundaries31. However, these nanostructured materials with much reduced κlat commonly exhibits reduction in their σ as well due to additional carrier scattering by either high-density grain boundaries or dislocations. This side effect of κlat reduction limits a further enhancement in ZT. To resolve this critical issue, incorporating multi-scale inclusions either intrinsically or extrinsically in a bulk TE matrix has been investigated because a band alignment between a matrix and an inclusion can circumvent a considerable reduction in σ by lowering an energy barrier in either conduction band or valence band between them.8, 32 However, such a TE bulk material has not been successfully realized for the Bi2Te3-based system due to the difficulty in forming well-distributed nano- and meso-sized inclusions, and the degradation of carrier mobility at interfaces between matrix and inclusions. To further decrease κlat of Bi2Te3-based materials, carefully chosen elements that can substitute into the Bi2Te3-based matrix are needed for additional phonon scattering from impurities. Especially, TE properties of transition metal (TM) doped Bi2Te3-based materials have not been well investigated, and additionally the d orbital electron of the TM may give a positive effect on TE performance by forming resonance levels.15 We tried to prepare Fe doped Bi0.4Sb1.6Te3 (BST) using a conventional high-temperature solid-state reaction in order 4


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to understand the effect of Fe doping on TE properties of BST. However, we found FeTe2 precipitates as a second phase in BST matrix because of solid solubility limit of Fe (< 1 mol. %) for Bi/Sb in BST, which was confirmed by formation energy calculation as well as X-ray diffraction measurement in this work. It has been also reported that the Bi0.48Sb1.50Fe0.02Te3 compound has a small amount of unidentified impurity peaks that should be related with FeTe2 phase.33 Thus, we synthesized BST-FeTe2 materials by the hightemperature solid-state reaction and then fabricated bulk pellets by melt-spinning (MS) and spark plasma sintering (SPS) in order to give a multi-scale structure including meso-scale grains, sub-micron sized FeTe2 inclusions along grain boundaries, and Fe-rich nanoinclusions within grains. In this work, we show that the BST-FeTe2 materials exhibit the enhanced ZT value due to the combined effect of high PF and low κlat. Very impressively, the highest ZT of 1.52 is achieved at 396 K for the BST-8 mol.% FeTe2, which is 43 % enhancement in ZT compared to the pristine BST (ZT = 1.06). We demonstrate that the Fe-rich nano-inclusions within grains as well as the sub-micron sized FeTe2 inclusions along grain boundaries provide an effective phonon scattering that results in a lower κlat than BST whereas the valence band alignment of BST and FeTe2 at the interface enables hole carriers to be easily transported without any energy barriers to overcome, leading to a high hole carrier mobility comparable to BST. This work proposes a guideline for designing high performance TE materials, suggesting that the interfacial electronic band alignment between a matrix and an inclusion as well as well-distributed nano- and meso-scale inclusions should be fulfilled.

RESULTS AND DISCUSSION 5



Nano- and micro-scale FeTe2 inclusions in the BST matrix are prepared by MS process under a high cooling rate (105~107 K/sec) to supersaturate the fourth element of Fe. The XRD patterns of SPS pellets of representative BST-x mol.% FeTe2 (x = 0, 4, 11) compositions are shown in Figure 1a. The XRD peaks from the samples are well matched with fundamental diffraction peaks from Bi0.4Sb1.6Te3 (JCPDS 65-3674, R-3m space group) and those from FeTe2 (JCPDS 14-0419, Pnnm space group) are also present when the nominal amount of FeTe2 is higher than 2%. Notably, the enlarged XRD patterns of BST-x% FeTe2 (x = 0, 2, 4, 8, 11) ranging between 30o and 35o in 2θ clearly show (111) and (120) peaks of FeTe2, indicating that the second phase of FeTe2 is successfully formed by a hightemperature solid state reaction followed by a subsequent rapid quenching by the MS process (Figure 1b). Furthermore, the chemical composition of SPS pellet analyzed by inductively coupled plasma atomic emission spectroscopy (ICP-AES) (Table S1) is almost same as nominal composition, indicative of a successful preparation of BST-FeTe2 materials. To further investigate the microstructure of the samples, an electron probe micro analysis (EPMA) is carried out to understand how FeTe2 phases are distributed on a microscale when their percentage increases and a transmission electron microscopy (TEM) study is especially performed in SPS pellets of BST and BST-8% FeTe2. The TEM image of BST-8% FeTe2 sample reveals that there exist various nanostructures including sub-micron sized second phase at the grain boundary (blue square in Figure 2a) and nanometer sized precipitates in the grain (yellow square in Figure 2a). Interestingly, the grain size of BST-8% FeTe2 ranges from 100 nm to a few µm, which is much smaller than that of BST (Figure S1a). However, it is noteworthy that the BST sample shows more complex bundle of dislocations compared to the BST-8% FeTe2 sample (Figure S1a), leading to the formation of cell structure with small angle grain boundaries as shown in Figure S1b. A high angle 6


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annular dark field (HAADF) scanning TEM (STEM) image shows an elemental sensitive contrast, revealing the distinct shape and size of the second phase and the precipitates (Figure 2b). The sub-micron sized second phase has an irregular particle shape and is identified to be FeTe2 phase in the selected area electron diffraction (SAED) patterns (inset of Figure

2a).

TEM-energy dispersive

X-ray spectroscopy

(EDS) elemental maps

corresponding to area of the HAADF-STEM image in Figure 2b shows that the inclusion is composed of Fe and Te, consistent with the XRD data in Figure 1, and its size distribution is diverse. When the amount of FeTe2 increases, the FeTe2 inclusions are more abundant, which is supported by EPMA images in (Figures S2 and S3). Furthermore, a large-scale microscopic investigation is also performed using a back scattered electron-scanning electron microscopy (BSE-SEM), indicating that there exist well-dispersed several micron-sized FeTe2 inclusions for BST-2% FeTe2 and BST-8% FeTe2 samples (Figure S4). This is in good agreement with both EPMA and low-magnification TEM analyses. Very interestingly, besides FeTe2 inclusions, we can easily find nanometer sized precipitates (< 20 nm in width) in the BST-8% FeTe2 sample as shown in the yellow square of Figure 2a. It appears that the precipitate grows along a van der Waals gap plane of BST lattice (Figures 2b and 3a). To elucidate the structural information of the nano-precipitates, the fast Fourier transformed (FFT) patterns are obtained from the matrix and the precipitate in Figure 3b. The FFT pattern of the matrix is in good agreement with the fundamental pattern of the BST phase, but that of the precipitate is not matched with the known patterns of either Fe or iron telluride phases. A TEM-EDS elemental map (Figure 3d) supports that it is a Ferich phase, but the identity of precipitate is unclear at this moment. The phase identification of the precipitate is in progress utilizing a three-dimensional atom probe tomography and we have a preliminary data, indicating that it may be a Fe oxide because it is supposed to have a 7



large lattice parameter according to the FFT pattern. One of the crystal plane is parallel to (001) plane of BST, but not the others. Consequently, the precipitate has a semi-coherent relationship with the BST matrix and there is a lattice mismatch between BST and the precipitate, resulting in formation of dislocations and thus significant strain around the precipitate as shown in Figures 3c and S5. We perform first-principles density functional theory (DFT) calculations to understand the effect of FeTe2 inclusions as well as Fe doping on TE properties of BSTFeTe2 materials. It has been well known that a doping mechanism in BST is determined through defects formation, dominantly by forming antisite and vacancy defects. Bi and Sb vacancies (VBi and VSb) generate holes while Te vacancies (VTe) donate electrons. In the case of antisites, either a Bi or Sb atom occupying a site of a Te atom (BiTe or SbTe) generates a hole. Figure 4a shows the formation energy (EFORM) of Fe in BST with various oxidation numbers (QOX) under Bi/Sb-rich (Te-poor) and FeTe2-rich condition. Our DFT calculations reveal that Fe prefers Bi/Sb site in BST (FeBi/Sb) showing a multi-valency nature, which is similar to a previous work34. The QOX of Fe is variable depending on the type of major charge carrier: either Fe4+ or Fe3+ is a stable valent state in p-type BST while Fe2+ is a stable state in n-type BST. In addition, the Fe has a meta-stable electronic spin configuration because the energy difference between high- and low-spin electronic configurations is only 30 meV. Here, we include the spin-orbit interaction (SOI) effect in the calculations and consequently the positions of band edges are lowered and the donor states (oxidation state of 4+) become more stable,35 compared to the calculation results without SOI. The multi-valent character of Fe implies that Fe is a charge carrier trap in BST, not acting as either an efficient acceptor or donor. As a result, Fe can trap holes and electrons under a p-type and n-type condition, respectively, indicating that a Fe doping in BST can decrease the p-type carrier concentration. 8


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The EFORM of FeBi/Sb defect is calculated to be 0.591 eV in p-type BST with the Fermi level being at the valence band maximum (VBM). In Figure 4b, we roughly estimate the solubility of Fe in BST because the Fe atomic density is simply predicted from the Boltzmann factor calculation with a formation energy of Fe and a synthesis temperature. When we assume that FeBi/Sb defects are generated at 800 K, the Fe atomic density is predicted to be 7 × 1017 cm-3 in BST compared to the BST atomic density of 4.33 × 1022 cm-3, indicating that the solid solubility limit of Fe for Bi/Sb in BST is about 0.4 %. Note that the measured carrier concentration of p-type BST in this work is 3.04 × 1019 cm-3 (Table 1) and thus a direct effect of Fe doping on the TE properties of BST should be negligible because of a very low Fe doping concentration in BST. Furthermore, to experimentally identify the valence state of Fe, we perform an X-ray photoelectron spectroscopy (XPS) analysis of the BST-x% FeTe2 samples. However, there exists only the XPS peak related with Fe-Te binding energy of FeTe2 phase36-37 and other XPS peaks of Fe-Te binding energy of (Bi,Sb,Fe)2Te3 are not observed, indicating that the Fe doping in the BST lattice is too small to be detected by XPS (Figure S6). Electrical transport properties of the samples in this work are measured perpendicular to the pressing direction of SPS unless otherwise noted. Figure 5a shows temperaturedependent electrical conductivity σ of MS-SPS BST-x% FeTe2 (x = 0, 2, 4, 8, 11). The σ of the samples monotonically decrease with increasing temperature, indicating a metallic conduction (degenerate semiconducting) behavior. Also, the σ follows a temperaturedependent power law of σ ~ T-β, and the β of the x = 0, 2, 4, 8, and 11 samples is 1.99, 1.63, 1.71, 1.68, and 1.65, respectively. The RT σ increases with increasing amount of FeTe2 until 4 % FeTe2 and then it drops with higher FeTe2 content (Table 1). To elucidate the origin of this RT σ trend, we performed Hall effect measurements of the samples. The hole carrier 9



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concentration np is calculated by the equation (1), using Hall coefficient RH measured by a Hall effect system: RH=1/npe

(1)

where e is an electron charge. The hole carrier mobility µp is obtained by the equation (2): σ = npeµp

(2)

The RT np of BST-x% FeTe2 (x = 0, 2, 4, 8, and 11) is tabulated in Table 1. It indicates that the introduction of FeTe2 phase in BST increases the RT np value up to a point, which is in good agreement with the variation of RT σ with FeTe2 content. However, Fe in BST could not be effective for p-type doping, as already elucidated from the DFT calculation in Figure 4. In addition, the change of RT np value with FeTe2 content is somewhat irregular, but the RT np values of the samples are repeatable within the error range of 5 %. We presume that the increased RT np value could be originated from a facile defect formation at the interface between BST and FeTe2 despite no consistent increase or decrease in the RT np with increasing FeTe2 content. The defect formation energy in either surface or interface is generally lower than the defect formation energy in bulk. The impurities also tend to be segregated near the interface. Thus, the RT np value could be enhanced with the incorporation of the sub-micron sized FeTe2 inclusions.34, 38-41 The temperature-dependent Seebeck coefficient S of the BST-x% FeTe2 (x = 0, 2, 4, 8, 11) is shown in Figure 5b. Very interestingly, the temperature variation of S of the samples is nearly identical to each other and the highest S values of 198 µV K-1 at RT and 219 µV K-1 at 420 K are achieved for BST-8% FeTe2, which are 5 – 7 % higher values than BST. Very impressively, in comparison with the BST sample, the change in S is significant compared to that in σ especially for the BST-8% FeTe2 sample (Table 1). However, the S value of the BST-11% FeTe2 is quite reduced compared to the BST-8% FeTe2. Considering that the RT np 10


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of BST-11% FeTe2 is similar to that of BST-8% FeTe2, the reduced S of the BST-11% FeTe2 could be caused by the enhancement on S in the conduction band, resulting in increasing the bipolar effect that is an interplay between band parameters of the valence band and those of the conduction band.42 To understand the high S value for FeTe2-incorporated BST, we do theoretical calculations including DFT and Boltzmann transport equation (BTE). The electronic band structure of FeTe2 is calculated within PBE+U (UFe = 4.0 eV)43 in Figure S7. Here we use marcasite phase of FeTe2 with the experimental lattice parameters (a = 5.27 Å, b = 6.27 Å, c = 3.87 Å),44 a non-parabolic band behavior being responsible for high TE performance is known for FeTe2,45 and the band gap of FeTe2 is reported to be larger than that of Bi2Te3. To overcome the band gap underestimation of FeTe2 in PBE and PBE+U calculations, we additionally apply the GOWO calculations46 and obtain the reliable band gap of 0.88 eV, consistent with the experimentally measured band gap of ~0.92 eV

47

. By

performing DFT calculations combined with BTE, we can estimate the S values of Bi0.4Sb1.6Te3 (Figure 5c). Here we use the full band structure in DFT to fully describe a nonparabolic character of BST and adopt the constant relaxation time approximation with the rigid band approximation. The S of Bi0.4Sb1.6Te3 is calculated by a method in a previous reference48. For comparison, we derive the carrier concentration-dependent S with various carrier effective mass m* including 1.2me, 1.4me, and 1.5me using a single parabolic band (SPB) model 49 based on the assumption of acoustic phonon scattering mechanism (scattering parameter r = -1/2). We find that the experimental data of BST-FeTe2 materials (red square symbols) are deviated from the Pisarenko line (black solid line) of BST in Figure 5c, revealing that there is an enhancement in S for BST-FeTe2 materials compared to BST (black square symbol), which should be attributed to the increase in effective mass. A very low Fe doping concentration in BST excludes a possible S enhancement mechanism due to Fe 11



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forming resonance levels in this work. There may be a modification of a scattering term (1/µ(E))(dµ(E)/dE) induced by an interfacial effect between BST matrix and FeTe2 inclusions in a following approximate formula (3):15 S=

గ మ ௞ಳ ଷ ௘

ሺ݇஻ ܶሻ ቂ

ଵ

ௗ௡ሺாሻ

௡ሺாሻ ௗா

+

ଵ

ௗఓሺாሻ

ఓሺாሻ ௗா

ቃ

(3)

We also calculate the S of FeTe2 and compare it with that of BST. It is surprising that the S value of FeTe2 is comparable to BST when the hole carrier concentration np is in the range of 1018 – 1020 cm-3. For np < 1018 cm-3, FeTe2 shows higher S values due to a larger band gap and smaller bipolar effect than BST. We believe that the high S value of FeTe2-incorporated BST should be attributed to the interfacial effect between BST matrix and FeTe2 inclusions. Figure 5d shows temperature-dependent power factor PF of BST-x% FeTe2 (x = 0, 2, 4, 8, 11). The FeTe2-incorporated samples show higher PFs than BST and especially their PFs are much higher at a high temperature ranging between 380 K and 480 K. The BST-8 % FeTe2 sample shows the highest RT PF value reaching 46.7 µW cm-1 K-2, which is enhanced by ~14 % compared to 41.0 µW cm-1 K-2 for BST. The RT µp of BST-x% FeTe2 (x = 0, 2, 4, 8, 11) is very unusual that even the incorporation of 11% FeTe2 in BST cannot significantly decrease the RT µp value of BST (Table 1). To understand the electrical transport properties of FeTe2-incorporated BST, we investigate the interfacial property of BST and FeTe2 by performing a vacuum energy alignment calculation for BST and FeTe2. The band diagram between BST and FeTe2 is shown in Figure 5e. The band edge positions of VBM and conduction band minimum (CBM) are obtained from the reference potential method50 with bulk and (0001) surface slab calculations and thus the band edge positions of BST are estimated by using the PBE with SOI calculations. Then the mid-gap energy, (EVBM + ECBM)/2, is corrected within GOWO 12


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approximation level without SOI. Finally, the band gap of BST is corrected by its experimental band gap48. The band edge positions of FeTe2 are fully obtained from the GOWO after PBE+U calculations. Strikingly, we find that there is an asymmetric band offset at the interface between Bi0.4Sb1.6Te3 and FeTe2. Interestingly, there is a large band gap difference between Bi0.4Sb1.6Te3 and FeTe2, but a transport barrier for hole charge carriers is very small (0.055 eV) whereas that for electron charge carriers is very large (0.61 eV). As a result, we suggest that the BST-FeTe2 composite can show a high p-type TE performance without a substantial degradation of µp value because of a very small barrier for hole charge carrier transport. Furthermore, we perform high-temperature Hall measurements of representative BST and BST-8% FeTe2 samples to understand their electrical transport scattering mechanisms. The np of BST and BST-8% FeTe2 monotonously increases with increasing temperature (Figure S8a), ranging from 3.04 × 1019 cm-3 and 3.54 × 1019 cm-3 at 300 K to 3.52 × 1019 cm-3 and 4.06 × 1019 cm-3 at 473 K, respectively. The µp follows a temperaturedependent power law of µp ~ T-λ, and the λ of BST and BST-8% FeTe2 is 2.16 and 1.93, respectively (Figure S8b). Thermal transport properties of the samples in this work are also measured along the same direction with electrical transport properties unless otherwise noted. Figure 5f show temperature-dependent thermal conductivity κ of the BST-x % FeTe2 (x = 0, 2, 4, 8, 11). All the samples show a typical temperature variation of κ for Bi2Te3-based materials, indicating that the κ decreases with increasing temperature until ~380 K and then increases with a further rise in temperature (bipolar conduction behavior). Furthermore, the κ of FeTe2incorporated samples show lower thermal conductivities than BST. The κ is the sum of the electrical thermal conductivity κel, which is evaluated by the Wiedemann-Franz relationship (κel = LσT, where L is Lorenz number), the bipolar thermal conductivity κbp, and the lattice 13



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thermal conductivity κlat. Based on the assumption of the SPB model with the acoustic phonon scattering,49 the Fermi integral can be estimated from the experimentally measured S value as the following equations (4) and (5): ܵ=±

௞ಳ ሺ௥ାହ/ଶሻிೝశయ/మ ሺకሻ ௘

൤ሺ௥ାଷ/ଶሻி

ೝశభ/మ ሺకሻ

ஶ

‫ܨ‬௡ ሺߦሻ = ‫׬‬଴

௫೙

ଵା௘ ሺೣష഍ሻ

− ߦ൨

(4)

݀‫ݔ‬

(5)

where ‫ܨ‬௡ ሺߦሻ and ߦ correspond to Fermi integral and reduced Fermi energy (EF-EC)/kT).

Then, the temperature variation of L can be calculated from the following equation (6): ௞

ଶ ሺ௥ା଻/ଶሻி

ሺకሻ

ሺ௥ାହ/ଶሻி

ሺకሻ

ଶ

ೝశఱ/మ ೝశయ/మ ‫ = ܮ‬ቀ ಳቁ ቈ − ൬ሺ௥ାଷ/ଶሻி ൰ ቉ ሺ௥ାଷ/ଶሻி ௘ ሺకሻ ሺకሻ ೝశభ/మ

ೝశభ/మ

(6)

where r = -1/2 for the acoustic phonon scattering. The temperature-dependent L of the samples are shown in Figure S9, which can be compared with the temperature-dependent L values extracted from Snyder’s method51 (Figure S10) that are similar to those from the SPB model. Taking the estimated L values of Figure S9, we calculated κlat + κbp (= κ – κel) of the BST-x% FeTe2 (x = 0, 2, 4, 8, 11) samples as shown in Figure 6a. Assuming that the κlat follows a temperature-dependent power law of κlat ~ T-1 dominated by a Umklapp phonon scattering, the temperature-dependent κlat of the samples can be approximately estimated (Figure S11). It is noteworthy that all the FeTe2-incorporated samples have lower κlat values than BST. For example, the RT κlat values of x = 0, 2, 4, 8 and 11 is 0.69, 0.56, 0.46, 0.44, and 0.56 W m-1 K-1, respectively, compared to the theoretical minimum lattice thermal conductivity (κmin = 0.31 W m-1 K-1) of BST calculated by the Cahill’s formula52-53. It should be noted that the RT κlat of the BST-11% FeTe2 is unexpectedly higher than that of BST-8% FeTe2, which should be attributed to compositional fluctuations, anisotropic crystal structure, or random distribution of inclusions with different sizes for nanostructured TE materials.54-56 14


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In addition, the size distribution and volume fraction of FeTe2 inclusions and Fe-rich precipitates should be critical to understand the phonon transport mechanism in the BSTFeTe2 composites, but it is technically very difficult to determine them because of a significantly different length scale between FeTe2 inclusions and Fe-rich precipitates. This result can be simply understood based on the microstructure of the BST-FeTe2 composites as discussed above, revealing that the nano- and meso-structures in BST-FeTe2 composites lead to significant drop in κlat originated from both boundary scattering and nanostructuring. Considering an elastic phonon scattering contribution, a Rayleigh scattering cross section can be approximately calculated to be ~ α6/λ4 based on the Rayleigh law, where α is the size of the particle and λ is the phonon wavelength.57 In general, atomic defects could only scatter short-wavelength phonons whereas particles larger than 5 nm could effectively scatter mid- to long-wavelength phonons. In this work, the κlat is estimated to be 0.42 W m-1 K-1 at 350 K in the BST-8% FeTe2 sample, indicative of a significant reduction in κlat by 38 % compared to the BST in this work. The combined effect of a smaller grain size, the sub-micron sized FeTe2 inclusions within grains, and Fe-rich nano-inclusions along grain boundaries followed by dislocations and strained area effectively scatter mid- to long-wavelength phonons, resulting in a considerable decease in κlat. Taking into account that the frequency-dependent electron mean free path of BST generally ranges from 1 to 100 nm,58 we believe that both the Fe-rich nano-inclusions and the sub-micron sized FeTe2 inclusions in BST-x% FeTe2 samples could negligibly scatter hole charge carriers, which is consistent with the trend of no considerable deterioration of µp at higher FeTe2 amount (Table 1). For better understanding the enhanced TE properties with FeTe2 incorporation, various plots of TE properties including σ, S, np, µp, and κlat at 300 K as a function of FeTe2 content are shown in Figure S12. 15



Dimensionless TE figure of merit ZT values of the samples in this work are obtained with the electrical and thermal transport properties being measured along the same direction. The ZT of the BST-x% FeTe2 (x = 0, 2, 4, 8, 11) samples are shown as a function of temperature (Figure 6b). After the introduction of FeTe2 in the BST matrix, we can observe a significantly enhanced ZT value than BST. The highest ZT of 1.52 is achieved at 350 ~ 400 K for 8% FeTe2-incorporated BST, which is enhanced by 43 % compared to ZT = 1.06 for BST. In addition, the temperature at which a maximum ZT is obtained is shifted up to the higher temperature range, which can be highly promising in the energy harvesting applications.

CONCLUSION In summary, we successfully fabricate tunable BST-FeTe2 heterostructures using a conventional high-temperature solid state reaction followed by MS-SPS method that could be a facile preparation method for a scale-up production. The heterostructures including the nano-precipitates of Fe-rich phase within grains and the sub-micron sized inclusions of FeTe2 at grain boundaries are prevalent in the BST-FeTe2 material, resulting in an effective boundary scattering of phonons. Furthermore, the valence band of FeTe2 are well aligned with that of BST so that hole charge carriers can easily pass across the interface between BST and FeTe2, which is elucidated by DFT calculation. Consequently, the p-type BST-8 mol.% FeTe2 material has a ZT > 1.5 in the temperature range of 350 – 400 K and thereby the BSTFeTe2 material having various nano-structures gives a great potential in the use for nextgeneration TE applications operating near and above room temperature.

EXPERIMENTAL SECTION 16


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Synthesis. Polycrystalline ingot samples of Bi0.4Sb1.6Te3–x mol.% FeTe2 (x = 0, 2, 4, 8, 11) compositions were prepared by melting the stoichiometric amount of high purity elements (>99.99 %) at 1373 K for 4 h and quenching in water. The ingots were put into a graphite nozzle with 0.4 mm diameter in an induction heater for a melt-spinning (MS) process and the rotating speed of Cu wheel (diameter ~250 mm) was fixed to 1000 rpm. Thin ribbons were produced by the MS process and the ribbons show amorphous structure on contact surface and crystalline nanostructure on free surface. The melt spun ribbons were mildly pulverized into powders and sintered using a spark plasma sintering (SPS) process at 673 K for 3 min under 60 MPa. Powder X-ray diffraction (XRD). Powder X-ray diffraction (XRD, New D8 Advance, Bruker, Cu Kα) was performed on the SPS pellets. Thermoelectric properties. The electrical conductivities (σ) and Seebeck coefficients (S) were simultaneously measured from room temperature to 473 K by a four-point probe method using the ZEM-3 apparatus (ULVAC-RIKO). The thermal diffusivities (D) were measured by a laser-flash analysis (LFA) using Netzsch LFA 457 and the thermal conductivity (κ) was calculated by the relation κ = dCPD, where d is the measured density of the sample and CP is the heat capacity measured by a differential scanning calorimeter (DSC). The room temperature d and CP values of the samples are shown in Table S2. Note that we measured the same sample multiple times for all the compositions, and especially the TE measurements were performed on two different batches of the BST-8% FeTe2 composition. For accurate estimation of ZT for the samples, we averaged the ZT values out for each composition.

17



Hall measurement. The Hall coefficient was measured by Hall effect measurement system (HT-Hall, ResiTest 8300, Toyo Corporation) to estimate carrier concentrations and carrier mobilities. Chemical analysis. Chemical compositions were determined by inductively coupled plasma atomic emission spectroscopy (ICP-AES) using Shimadzu ICP 8100. X-ray photoelectron spectroscopy (XPS). XPS spectra of the samples were measured by a VersaProbe PHI 5000 (ULVAC-PHI). Transmission electron microscopy (TEM). The thin foil TEM sample was prepared by mechanical polishing followed by Ar ion milling (PIPS, Gatan Corporation). For the microstructure analysis, scanning transmission electron microscopy (STEM, Titan 80-300, FEI Corporation) equipped with Cs probe corrector. The elemental composition was investigated using TEM (Talos, FEI Corporation) equipped with super X- energy dispersive spectroscopy (EDS). Electron probe micro analysis (EPMA). The size and distribution of precipitates were characterized by using EPMA (JXA 8530F, JEOL). Back scattered electron-scanning electron microscopy (BSE-SEM). The size and distribution of precipitates were characterized by using a FEI Nova NanoSEM 450. Density functional theory (DFT) calculations. First principles DFT calculations59-60 are performed to understand the effect of FeTe2 inclusion on the defect and thermoelectric properties of BST. Here, we used a planewave basis set, projector-augmented-wave (PAW) pseudopotentials,61 and a generalized gradient approximation parameterized by PerdewBurke-Ernzerhof (PBE),62 which are implemented in VASP code63-64. For the atomic model 18


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for BST, we consider a (4×4×1) hexagonal supercell, containing 24 Bi, 72 Sb, and 144 Te atoms, where the atomic positions of Bi and Sb are assumed to be ordered in BST. The lattice parameters of BST are interpolated from the experimental values from Bi2Te3 and Sb2Te3 (a = 4.2767 Å, c = 30.4174 Å), considering the lattice parameter from Bi0.4Sb1.6Te3. The atomic positions are fully relaxed until the remaining forces less than 0.008 eV/Å. The electronic structures are calculated for optimized structure with a inclusion of spin-orbit-interaction, which is found to be very important to describe the band structure and the formation energy of defects for heavy element containing materials.65 The formation energy EFORM of defect is calculated by using the formalism in the literature.66-68 Note that there is a difficulty in predicting the band structure of Bi2Te3 and the shallow defects with hybrid functionals.69-70

ASSOCIATED CONTENT Supporting Information Data of additional TEM images, EPMA images, XPS, electronic band structure of FeTe2, temperature-dependent hole concentration and mobility, temperature-dependent Lorenz number, temperature-dependent lattice thermal conductivity, and tables of chemical composition, room temperature density and specific heat values. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION Corresponding Author 19



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*[email protected] Author Contributions #

W.H.S. and J.W.R. equally contributed to this work.

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT W.H.S., S.L., and W.S.S. were supported by a grant from the Fundamental R&D program for Core Technology of Materials (10048035) funded by the Ministry of Trade, Industry, and Energy

(MOTIE)

and

by

National

Research

Foundation

of

Korea

(NRF-

2017R1D1A1B03034322). B.R. was supported by the Korea Institute of Energy Technology Evaluation and Planning (KETEP) and MOTIE of the Republic of Korea (No. 20162000000910). H.J.C. was supported by the Korea Institute of Science and Technology (KIST) Institutional Program (2V05210).

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Figure Legends Figure 1. Powder XRD patterns of (a) SPS pellets of BST, BST-4% FeTe2, and BST-11% FeTe2 together with theoretical patterns of Bi0.4Sb1.6Te3 and FeTe2 in the 2θ range of 10 – 60°, 28


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and (b) BST, BST-2% FeTe2, BST-4% FeTe2, BST-8% FeTe2, and BST-11% FeTe2 in the 2θ range of 30 – 35°. Figure 2. (a) Low-magnification TEM image of the SPS pellet of BST-8% FeTe2 and an inset SADP corresponding to the second phase, FeTe2 inclusion (blue square region), (b) HAADF-STEM image showing various nanostructures in BST-8% FeTe2, and (c) TEM-EDS elemental maps of Bi, Sb, Te, and Fe, corresponding to HAADF STEM image in (b). Figure 3. (a) High resolution HAADF-STEM image of the nanoscale precipitate. (b) FFT patterns from the matrix (red square region) and the precipitate (blue square region), respectively. (c) Filtered high resolution HAADF-STEM image showing existence of dislocations near the interface between the matrix and the nano-precipitate, resulting from the lattice mismatch. (d) HAADF-STEM image and corresponding TEM-EDS elemental maps of Bi, Sb, Te, and Fe of the precipitate. Figure 4. (a) Formation energy of FeBi/Sb (EFORM) in BST is calculated as a function of Fermi level (EF) with various QOX values of Fe. The different line slope represents different QOX. (b) The Fe atomic density in BST is estimated based on the Boltzmann Factor with formation energy and synthesis temperature. The dotted vertical and horizontal line represents the melting temperature of Bi2Te3 (Tm) and the hole concentration of 4×1019 cm-3, respectively. Figure 5. Thermoelectric properties taken perpendicular to the pressing direction of SPS as a function of temperature for BST, BST-2% FeTe2, BST-4% FeTe2, BST-8% FeTe2, and BST11% FeTe2. (a) Electrical conductivity σ. (b) Seebeck coefficient S. (c) The estimated room temperature S of BST and FeTe2 are calculated as a function of hole carrier concentration by a Full band BTE and a SPB model (r = -1/2). The experimental data in this work are plotted as black open square (BST) and red open square symbols (BST-FeTe2). (d) Power factor PF 29



(= σS2). (e) The predicted electronic band diagram between BST and FeTe2. Note that there is negligible band offset for valence band maximum (VBM), which could enable the hole carrier to easily transport across the interface. f) Total thermal conductivity κ. Figure 6. Thermoelectric properties taken perpendicular to the pressing direction of SPS as a function of temperature for BST, BST-2% FeTe2, BST-4% FeTe2, BST-8% FeTe2, and BST11% FeTe2. (a) Lattice thermal conductivity plus bipolar thermal conductivity κlat + κbp. (b) Thermoelectric figure of merit ZT.

30


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Figure 1. Powder XRD patterns of (a) SPS pellets of BST, BST-4% FeTe2, and BST-11% FeTe2 together with theoretical patterns of Bi0.4Sb1.6Te3 and FeTe2 in the 2θ range of 10 – 60°, and (b) BST, BST-2% FeTe2, BST-4% FeTe2, BST-8% FeTe2, and BST-11% FeTe2 in the 2θ range of 30 – 35°.

31



Figure 2. (a) Low-magnification TEM image of the SPS pellet of BST-8% FeTe2 and an inset SADP corresponding to the second phase, FeTe2 inclusion (blue square region), (b) HAADFSTEM image showing various nanostructures in BST-8% FeTe2, and (c) TEM-EDS elemental maps of Bi, Sb, Te, and Fe, corresponding to HAADF STEM image in (b).

32


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Figure 3. (a) High resolution HAADF-STEM image of the nanoscale precipitate. (b) FFT patterns from the matrix (red square region) and the precipitate (blue square region), respectively. (c) Filtered high resolution HAADF-STEM image showing existence of dislocations near the interface between the matrix and the nano-precipitate, resulting from the lattice mismatch. (d) HAADF-STEM image and corresponding TEM-EDS elemental maps of Bi, Sb, Te, and Fe of the precipitate.

33



Figure 4. (a) Formation energy of FeBi/Sb (EFORM) in BST is calculated as a function of Fermi level (EF) with various QOX values of Fe. The different line slope represents different QOX. (b) The Fe atomic density in BST is estimated based on the Boltzmann Factor with formation energy and synthesis temperature. The dotted vertical and horizontal line represents the melting temperature of Bi2Te3 (Tm) and the hole concentration of 4×1019 cm-3, respectively.

34


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Figure 5. Thermoelectric properties taken perpendicular to the pressing direction of SPS as a function of temperature for BST, BST-2% FeTe2, BST-4% FeTe2, BST-8% FeTe2, and BST11% FeTe2. (a) Electrical conductivity σ. (b) Seebeck coefficient S. (c) The estimated room temperature S of BST and FeTe2 are calculated as a function of hole carrier concentration by a Full band BTE and a SPB model (r = -1/2). The experimental data in this work are plotted as black open square (BST) and red open square symbols (BST-FeTe2). (d) Power factor PF (= σS2). (e) The predicted electronic band diagram between BST and FeTe2. Note that there is negligible band offset for valence band maximum (VBM), which could enable the hole carrier to easily transport across the interface. f) Total thermal conductivity κ.

35



Figure 6. Thermoelectric properties taken perpendicular to the pressing direction of SPS as a function of temperature for BST, BST-2% FeTe2, BST-4% FeTe2, BST-8% FeTe2, and BST11% FeTe2. (a) Lattice thermal conductivity plus bipolar thermal conductivity κlat + κbp. (b) Thermoelectric figure of merit ZT.

36


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Table 1. Room temperature σ, S, np, and µp of BST, BST-2% FeTe2, BST-4% FeTe2, BST8% FeTe2, and BST-11% FeTe2. σ (S cm-1) at 300 K

S (µV K-1) at 300 K

np (1019 cm-3) at 300 K

µp (cm2 V-1 s-1) at 300 K

Bi0.4Sb1.6Te3 (BST)

1199

185

3.04

247

BST-2% FeTe2

1376

172

4.24

203

BST-4% FeTe2

1379

173

4.36

198

BST-8% FeTe2

1192

198

3.54

210

BST-11% FeTe2

1135

181

3.68

193

37



TOC:

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Enhancing Thermoelectric Performances of Bismuth Antimony

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