Enhancing p-type thermoelectric performances of polycrystalline SnSe

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Enhancing p-type thermoelectric performances of polycrystalline SnSe via tuning phase transition temperature Yong Kyu Lee, Kyunghan Ahn, Joonil Cha, Chongjian Zhou, Hyo Seok Kim, Garam Choi, Sue In Chae, Jae-Hyuk Park, Yeseul Lee, Cheol-Hee Park, Sung-Pyo Cho, Sang Hyun Park, Yung-Eun Sung, Won Bo Lee, Taeghwan Hyeon, and In Chung J. Am. Chem. Soc., Just Accepted Manuscript • DOI: 10.1021/jacs.7b05881 • Publication Date (Web): 14 Jul 2017 Downloaded from http://pubs.acs.org on July 14, 2017

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Enhancing p-type thermoelectric performances of polycrystalline SnSe via tuning phase transition temperature Yong Kyu Lee,†,‡ Kyunghan Ahn,‡ Joonil Cha,†,‡ Chongjian Zhou,‡ Hyo Seok Kim,‡ Garam Choi, ‡ Sue In Chae,†,‡ Jae-Hyuk Park,†,‡ Yeseul Lee,⊥ Cheol-Hee Park,⊥ Sung-Pyo Cho,§ Sang Hyun Park,∥ Yung-Eun Sung,†,‡ Won Bo Lee,‡ Taeghwan Hyeon,†,‡ In Chung *,†,‡

†

Center for Nanoparticle Research, Institute for Basic Science (IBS), Seoul 08826, Republic of Korea. School of Chemical and Biological Engineering and Institute of Chemical Processes, and §National Center for Inter-University Research Facilities, Seoul National University, Seoul 08826, Republic of Korea. ‡

⊥

LG Chem/Research Park, Daejeon 34122, Republic of Korea. Advanced Materials and Devices Laboratory, Korea Institute of Energy Research, Daejeon 34129, Republic of Korea.

∥

* To whom correspondence should be addressed: [email protected] ABSTRACT SnSe emerges as a new class of thermoelectric materials since recent discovery of an ultrahigh thermoelectric figure of merit in its single crystals. Achieving such performance in the polycrystalline counterpart is still challenging and requires fundamental understandings of its electrical and thermal transport properties as well as structural chemistry. Here we demonstrate a new strategy of improving conversion efficiency of bulk polycrystalline SnSe thermoelectrics. We show that PbSe alloying decreases the transition temperature between Pnma and Cmcm phases and thereby can serve as a means of controlling its onset temperature. Along with 1% Na doping, delicate control of the alloying fraction markedly enhances electrical conductivity by earlier initiation of bipolar conduction while reduces lattice thermal conductivity by alloy and point defect scattering simultaneously. As a result, a remarkably high peak ZT of ~ 1.2 at 773 K as well as average ZT of ~ 0.5 from RT to 773 K is achieved for Na0.01(Sn1-xPbx)0.99Se. Surprisingly, spherical-aberration corrected scanning transmission electron microscopic studies reveal that NaySn1-xPbxSe (0 < x  0.2; y = 0, 0.1) alloys spontaneously form nanoscale particles with a typical size of ~ 5 – 10 nm embedded inside the bulk matrix, rather than solid solutions as previously believed. This unexpected feature results in further reduction in their lattice thermal conductivity. sixty years.18-30 Because of scantiness and high cost of Te element, high performance TE materials with no Te have been highly sought after. SnSe that consists of no scarce and toxic elements has been ignored for TE applications mainly because of seemingly high electrical resistivity. In 2014, SnSe single crystal was reported to show a record high ZT of 2.6 at 923 K along the crystallographic b-axis. Its exceptionally high ZT originates from unique crystal chemistry.31 The corrugated layered structure coupled with highly distorted [SnSe7] polyhedra provides giant anisotropic and anharmonic bonding characteristic, consequently giving rise to one of the lowest κlatt values known for crystalline materials (< 0.4 W m-1 K-1 at 923 K).32-36 Another key element is a displacive phase transition from the Pnma to Cmcm space group at 750 – 800 K.31,37,38 This structural change causes the considerably reduced band gap (0.61 eV to 0.39 eV by theoretical calculations)31 and resultantly increases

INTRODUCTION Thermoelectric (TE) power generation can directly convert heat into electricity, thereby a promising renewable energy technology for a wide range of applications including flexible electronics, automobiles, and industrial power plants.1-11 Its characteristic advantages are noiseless operation and high mechanical reliability due to a device structure with no moving parts.12-17 The prime interest of this technology is further improving TE conversion efficiency for broader commercial markets. The efficiency of a TE material is determined by the dimensionless TE figure of merit ZT = S2σT/κ, where S is the Seebeck coefficient, σ is the electrical conductivity, κ is the thermal conductivity, T is the absolute temperature, and S2σ is power factor (PF). Bismuth telluride and lead telluride have served as top performing TE materials for the past 1

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conductivity of the materials.

carrier mobility and electrical conductivity due to a bipolar conduction process. As a consequence, SnSe features an abrupt rise of ZT after undergoing phase transition. Note that synthesizing sizable, qualifying single crystals is the expensive, lengthy, and laborintensive processes, and hence they are unsuitable for mass production and practical thermoelectric applications. However, polycrystalline SnSe exhibits inferior performance to its single crystalline counterpart yet. In contrast to Pb chalcogenides, effective chemical dopants to SnSe are limited.39 Ptype dopants of Ag, Na, and K and n-type dopants of I and BiCl3 were introduced and enhanced a ZT value.39-44 On the other hand, hydrothermal synthesis to incorporate PbSe to SnSe yields phase-separated nanopowders.45 PbSe alloying up to 12 mol% with Na0.01Sn0.99Se by solid-state reaction was reported to form solid solutions but deteriorates ZT from 0.85 of the pristine to less than ~ 0.7 of the best performing alloy,46 which is in contrary to the results described in the work. However, deep understandings of structural features at the atomic level as well as electrical and thermal transport, and consequent guidelines for further enhancing its TE performance are highly requested.

EXPERIMENTAL SECTIONS Reagents. The following reagents were used as obtained unless note otherwise: Sn chunk (99.999% American Elements, US), Se shot (99.999% 5N Plus, Canada), Na pieces (99.9%, Sigma-Aldrich, US), and Pb wire (99.99%, 5N Plus, Canada). Synthesis. To synthesize polycrystalline Sn1-xPbxSe (x = 0, 0.03, 0.04, 0.05, 0.06, 0.10, 0.15, 0.20) ingots and those doped with 1% Na, a stoichiometric mixture of constituent elements were loaded to an carbon-coated, evacuated fused silica tube (~ 10–4 Torr). The total weight of starting materials were typically ~15 g. The reaction tube was inserted into a larger evacuated tube to protect the material from oxidation in the case of cracks of the inner tube due to the phase transition during a cooling process. The double sealed tube was heated to 1223 K over 12 h and soaked at that temperature for 12 h, followed by a furnace cooling to room temperature. We found that a ZT was lowered when the reaction tube was quenched to air. The obtained ingot was powdered by a mechanical grinder, sieved to < 45 μm, and loaded into a graphite die in an Ar-filled glove box. The powders were densified at ~ 783 K for 5 min under an axial pressure of 50 MPa in a vacuum of ~ 1.4  10−2 Torr using spark plasma sintering (SPS) (SPS-211Lx, Fuji Electronic Industrial Co., Japan). Typical images of SPS-processed pellets and their dimensions are given in Figure S1. For the samples after SPS, chemical compositions were characterized by field emission electron probe microanalyzer (EPMA) using a JEOL JXA-8530FPlus HyperProbe field emission EPMA and inductively coupled plasma atomic emission spectroscopy (ICPAES) using a PerkinElmer Optima 8300. The obtained compositions are close to the nominal ones (Table S1).

Here we report a new strategy to improve TE properties of bulk polycrystalline SnSe, namely, tuning the temperature of phase transition (Tc) by impurity alloying. We demonstrate that PbSe alloying to form Sn1-xPbxSe facilitates the Pmna-Cmcm phase transition and reduces the Tc at a rate of ~ – 3.2 K per x (0 < x  0.2). For example, the Tc for SnSe at 795 K significantly drops to 730 K for Sn0.8Pb0.2Se. Very importantly, the decreased Tc lowers the onset temperature of bipolar conduction occurring at high temperatures and consequently both σ and κ start to rise above the Tc. To be maximally effective PbSe alloying, we also discover that Na doping at the optimal level is essential. It exclusively enhances σ due to boosting carrier concentration along with suppression of κ by point defect scattering. As a result, p-type Na0.01(Sn0.95Pb0.05)0.99Se sample exhibits an average ZT of ~ 0.5 from RT to 773 K and a maximum ZT of ~ 1.2 at 773 K. The achieved values are the highest among bulk-processed polycrystalline SnSebased materials. The latter value is two-fold magnitude larger than that of undoped SnSe at 823 K.40 The enhanced ZT is attributed to the synergistic effect of an enhanced PF by the decreased Tc and a reduced lattice κ by phonon scattering due to chemical doping and alloying. Employing atomic resolution scanning transmission electron microscope, we reveal that PbSe alloying generates a substantial degree of nanoscale precipitates rather than forms solid solutions, in contrary to the published phase diagram47 and literature,46 which further reduces lattice thermal

Powder X-ray diffraction (XRD). Powder XRD analysis was performed using a SmartLab Rigaku powder X-ray diffractometer (Cu Kα graphitemonochromatized radiation) operating at 40 kV and 30 mA at room temperature. In situ temperaturedependent powder XRD patterns were collected from 740 K to 830 K at a rate of 10 K h–1 on a SmartLab Rigaku powder X-ray diffractometer with high temperature accessory operating at 45 kV and 200 mA. Differential scanning calorimeter (DSC) analysis. To detect phase transition, the DSC measurement was performed on a Netzsch DSC 214 Polyma instrument in an aluminum (99.5 %) crucible with a lid under Ar flow to 550 K at a heating rate of 10 K min-1. Electrical and thermal transport property measurements. The samples after SPS were cut and 2



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GATAN). For STEM-EDS analyses, chemical maps were taken with a probe size of 0.13 nm and a probe current of 40 pA. For STEM-EELS measurements, the energy dispersion was set at 0.1 eV ch−1. The fullwidth at half-maximum of the zero-loss peak in vacuum was 0.8 eV. The convergence and collection semiangles were 19 and 19.8 mrad, respectively.

polished into a disk with a radius of 13 mm and thickness of ~ 2 mm under a N2 atmosphere. The electrical conductivity and Seebeck coefficient were measured simultaneously under an Ar atmosphere from room temperature to 823 K using a Netzsch SBA 458 Nemesis system. Temperature-dependent Hall coefficients (RH) of the samples were measured by a Lake Shore HMS8407 Hall effect measurement system in a magnetic field of 1.5 T. The hole carrier concentration (nH) and hole mobility (μH) were calculated by nH = 1/(e∙RH) and μH = RH∙σ, respectively. A Netzsch LFA 457 MicroFlash instrument was used to measure the thermal diffusivity of the samples coated with graphite. The thermal conductivity was calculated from κtot = D∙Cp∙ρ, where D is the thermal diffusivity, Cp is the heat capacity, and ρ is the mass density of the specimens. The temperature-dependent Cp values were indirectly derived using a standard sample (Pyroceram).28,31 The ρ values used were obtained using their geometrical dimensions and masses (Table S2). The total thermal conductivity κtot is the sum of the lattice (κlatt) and electronic thermal (κelec) conductivities. κelec is proportional to the electrical conductivity (σ) according to the Wiedemann−Franz law (κelec = L∙ σ ∙T), where L is the temperature-dependent Lorenz number and T is the absolute temperature. An L value as a function of temperature were obtained from previous studies.31 The κlatt value was estimated by subtracting the κelec value from the κtot value: κlatt = κtot – κelec.

Electronic structure calculations. First principles electronic structure calculations were performed within the density functional theory (DFT) scheme. The projector augmented wave method48 with the plane-wave basis set was employed, implemented in Vienna ab initio simulation package.49 Within the generalized gradient approximation, the PerdewWang form of the exchange-correlation functional50 was used. A plane-wave energy cutoff of 400 eV was used with the 6  6  6 meshes of special k-points in the first Brillouin zone. The electronic convergence and force convergence were specified as 10-7 eV and 10-4 eV Å-1. To describe PbSe alloying in SnSe, we used 1  2  2 supercell of the conventional SnSe unit cell, which accommodates 32 atoms. For PbSe alloying, one Pb atom substitutes for one Sn atom. According to our electronic structure calculations on SnSe (Sn4Se4) and Sn3Pb1Se4 both with and without spin-orbit coupling, the difference in band gaps were nearly negligible, consistent with the previous report.51 RESULTS AND DISCUSSION

Transmission electron microscopy. Cross-sectional samples for scanning TEM (STEM) were prepared by focused ion beams (FIB, Helios 650, FEG, FEI) with a dual beam microscope using gallium ion milling. Before the ion milling process, carbon was sputtered to preserve the sample by surface coating. The specimens were further polished with a low-voltage and low-angle argon ion beam milling apparatus (NANO MILL, Model 1040, FISCHIONE). The atomic structure and chemical composition were investigated employing a spherical aberrationcorrected JEM ARM-200F microscope (Cold FEG Type, JEOL) equipped with a SDD type energy dispersive X-ray spectroscopy (EDS) detector (Solid Angle 0.9-sr, X-MaxN 100TLE, OXFORD) and electron energy loss spectroscopy (EELS) detector (965 GIF Quantum ER, GATAN) at 200 kV, which are installed at the National Center for Inter-university Research Facilities (NCIRF) at Seoul National University. In the high-angle annular dark-field (HAADF) STEM images, the point-to-point resolution was approximately 80 pm after the Cscorrection, and the angular range of the annular detector used was 68 mrad to 280 mrad. All images were recorded by a high-resolution CCD detector using a 2 k × 2 k pixel device (UltraScan 1000,

Tuning Phase Transition Temperature of SnSe via PbSe alloying Chalcogen salts of lead (PbQ; Q = S, Se, Te) adopt the cubic rock-salt structure. In contrast, those of tin (SnQ), its smaller congener in Group 14, offer different crystal chemistry. SnSe and SnS crystallize in the orthorhombic GeS-type structure consisting of corrugated two-atom-thick SnSe layers, whereas SnTe adopts the same structure as the Pb salts. As a result, their alloys of Sn1-xPbxSe and Na-doped phases of Na0.01(Sn1-xPbx)0.99Se form the GeS-type structure only up to x ~ 0.2 according to powder X-ray diffraction (PXRD) patterns (Figures 1a and 1b), in contrary to a continuous solid solution of Sn1-xPbxTe (0 x 1) composed of isostructural SnTe and PbTe.52 The published bulk phase diagram of PbSe-SnSe system 47,53,54 supports the PXRD results (Figure 2). Accordingly, the higher degree of PbSe alloying (x 0.2) should result in segregation of PbSe as observed in the sample of x = 0.3 (Figure 1a). The refinement of the XRD patterns for Sn1-xPbxSe and Na0.01(Sn1xPbx)0.99Se reveals the lattice expansion along the aand b- axes and the contraction along the c-direction upon incorporation of larger Pb and Na atoms than Sn 3


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element to exceptional thermoelectric (TE) performance of SnSe is a temperature-driven displacive phase transition, we performed in situ temperature-dependent PXRD studies for pristine SnSe, Sn0.95Pb0.05Se, and Sn0.9Pb0.1Se (Figure 3) on heating from 740 K to 830 K at a rate of 10 K h-1. The characteristic (111) reflection peak of the Pnma at 2θ = 30.1˚ gradually disappears from 740 to 830 K in Sn0.9Pb0.1Se while it still keeps intensity in pristine SnSe. This observation demonstrates that Pb alloying effectively reduces the onset temperature of phase transition (Tc) of SnSe, which is further confirmed by differential scanning calorimetry (DSC) studies.

Figure 1. Powder XRD patterns of (a) Sn1-xPbxSe (x = 0 – 0.3) and (b) Na0.01(Sn1-xPbx)0.99Se (x = 0 – 0.2) materials with standard diffraction patterns of SnSe (The International Centre for Diffraction Data (ICDD) PDF 01-075-6133) and PbSe (00-006-0354). PbSe Bragg peaks are marked by asterisks.

Figure 3. In situ temperature-dependent PXRD patterns of (a) SnSe, (b) Sn0.95Pb0.05Se, and (c) Sn0.9Pb0.1Se heated from 740 K to 830 K with the interval of 10 K, showing temperature-driven phase transition from the Pnma to Cmcm space group. The (111) reflection peak of the Pnma is gradually weakened on heating. DSC analysis of Sn1-xPbxSe reveal that the Tc decreases almost linearly at a rate of ~ – 3.2 K per x with increasing x up to ~ 0.2, the solubility limit of Pb (Figure 4). For example, Sn0.8Pb0.2Se undergoes the transition at 730 K, which is 65 K lower than that of pristine SnSe at 795 K. This trend can be understood in light of the relationship between defect formation and change of the Gibbs free energy (G = H – TS) that determines the phase transition phenomenon. PbSe alloying can be regarded as forming the defect in SnSe crystal structure and would be endothermic because the resulting lattice disruption raises the enthalpy (H) of Sn1-xPbxSe (Figure 4). At the same time, the induced disorder also contributes to elevating its entropy (S). The –TS term becomes more negative with higher PbSe content or temperature. Upon heating, the minimum in the Gibbs energy shifts to lower temperature with the higher PbSe alloying fraction. Correspondingly, the Tc of Sn1-xPbxSe could diminish with increasing x. This finding is substantially important in SnSe thermoelectrics exhibiting the considerable decrease in bandgap (~ 0.22 eV by theoretical calculations) brought by the Pnma (0.61 eV) to Cmcm (0.39 eV) transition.31 In this case, a

Figure 2. A bulk phase diagram of the PbSe-SnSe ternary system adapted from ref 47. atom (Figure S2). Note that the effective ionic radius of Pb2+, Na+, and Sn2+ in octahedral geometry is 1.19, 1.02, and 0.87 Å, respectively.55,56 For example, the lattice dimensions of Sn0.8Pb0.2Se are a = 11.601(6) Å, b = 4.197(1) Å, c = 4.424(1) Å, and V = 215.402(9) Å3 and those of SnSe are a =11.500(2) Å, b =4.158(2) Å, c =4.447(1) Å, and V = 212.642(8) Å3. Because a key 4



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bipolar conduction process occurs with rising temperature, potentially leading to the increase in electrical and thermal conductivity.31 This implies that tuning the Tc could be a means of enhancing the ZT of SnSe thermoelectrics.

phases negligibly deviate to be at 0.81 ± 0.04 eV and 0.41 ± 0.04 eV with a wide range of Pb concentrations, respectively, which is consistent with the experimental results (Figure S6).

Figure 4. Differential scanning calorimetry analysis of Sn1-xPbxSe (x = 0 – 0.2) at a rate of 10 K min-1 on heating, showing an endothermic peak for temperature-driven phase transition.

Thermoelectric properties of PbSe alloyed SnSe

Comparison of the projected density of states for SnSe and Sn0.8125Pb0.1875Se explains why Pb substitution for Sn does not affect the magnitude of energy gaps (Figure S7). According to our calculations for SnSe, the VBM and the CBM mainly consist of p orbitals of Se and Sn, respectively. As a consequence, the band gap originates from the p-p mixing between Se and Sn. Contribution from p orbitals of Pb is shown only inside the conduction band (~ – 8.54 eV – – 0.07 eV) rather than near the Fermi level with a similar band dispersion to that of Sn, thereby negligibly changing the CBM and the magnitude of the energy gaps of Sn1-xPbxSe materials.

The solid-state optical absorption spectra of Sn1-xPbxSe (0 ≤ x ≤ 0.2) reveal an unusual trend in their energy gaps as a function of Pb concentration (Figure S3). The energy gaps keep constant at 0.86 eV despite the substantial molar fraction of Pb up to ~ 20% and the covalent nature of the Pb-Se bonding. It is striking contrast to other ternary alloys such as Cd1xHgxTe and Pb1-xSnxTe. The former exhibits a linear dependence of band energies as a function of Hg concentration57 and the latter, a rare case, shows an anomalous band inversion phenomenon upon Sn substitution.58,59,60 The origin of this observation is discussed below in conjunction with the electronic structure calculations. Electronic Structure Calculations To better understand the experimental findings and their origin, we performed ab initio density functional theory calculations for both low (Pnma) and high temperature phases (Cmcm) of SnSe and Sn1-xPbxSe (x = 0.0625, 0.125, 0.1875). For both phases, PbSe alloying does not noticeably disturb their electronic structures (Figures S4 and S5). For the Pnma structures, the valence band maximum (VBM) and the conduction band minimum (CBM) occur along the Γ-Z and Γ-Y direction, respectively. For the Cmcm structures, the VBM and the CBM are shown along the Γ-Y and the Γ-Z direction, respectively. Resultantly, their indirect bad gap nature is confirmed. The theoretical energy gaps of the Pnma and Cmcm

To verify our strategy of enhancing ZT via controlling the Tc, we investigated the Sn1-xPbxSe system for the relationship between the Tc and TE performance (Figure 5). All samples were synthesized in ingot form by melting constituent elements. They were powder-processed and spark plasma sintered into dense pellets. Because of the layered structure of SnSe and the characteristic of SPS process, TE properties of Sn1-xPbxSe and its Na-doped systems are highly anisotropic. We measured electronic and thermal transport properties both perpendicular and parallel to the press direction. Because ZTs along the former direction is higher than that along the latter, we mainly discuss TE properties of the materials developed in this work taken along the former direction, if noted otherwise. Their TE properties obtained along the latter direction are presented in Supporting Information (Figure S12). The temperature-dependent electrical conductivity (σ) of Sn1-xPbxSe keeps almost constant, very low values less than 2 S cm–1 up to ~ 300 – 523 K and abruptly rises afterwards (Figure 5a). The upturn of the σ mainly results from the thermal excitation of minority carriers that is strongly related with the Pnma-Cmcm phase transition.31,61 The σ value at high temperature decreases gradually with the higher degree of PbSe alloying because of the reduced hole concentration (nH) and mobility (μH), as confirmed by the Hall effect measurement. The nH value sharply drops from the order of 1017 cm-3 for x ≤ 0.1 to 1016 cm-3 for x > 0.1 at 300 K. Plausibly, PbSe alloying considerably suppresses the generation of intrinsic Sn vacancies (Figure S8a). The μH value lessens from 46 cm2 V-1 s1 for SnSe to 11 cm2 V-1 s-1 for Sn0.8Pb0.2Se because of

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Figure 5. Temperature dependence of (a) electrical conductivity, (b) Seebeck coefficient, (c) power factor, (d) total thermal conductivity, (e) lattice thermal conductivity, and (f) figure of merit ZT of Sn1-xPbxSe (x = 0 – 0.2). enhanced point defects and alloy scattering with increasing Pb amounts (Figure S8b).

temperatures. A peak TE figure of merit ZT for Sn1-xPbxSe declines with PbSe alloying from 0.68 for pristine SnSe to 0.44 for Sn0.8Pb0.2Se (Figure 5f). We were able to tune the Tc by PbSe alloying, which directly shifts the onset temperature of bipolar conduction. Bipolar conduction increases both σ and κ at high temperatures. However, σ is so poor that κ is dominantly reflected on ZT of Sn1-xPbxSe. Note that Sn4+ is ubiquitous even in Sn+2 compounds, accounting for their p-type conduction behavior, in contrast to Pb almost exclusively adopting the formal charge of 2+. As a result, PbSe alloying results in the 1017 cm-3 and the low carrier density of ~ 1016 resultant poor σ of the materials. As a consequence, boosting hole concentration of Sn1-xPbxSe system up to the order of ~ 1019 cm-3 is essential to benefit tuning effect of the Tc by impurity alloying.

The temperature-dependent Seebeck coefficient (S) of Sn1-xPbxSe sharply decreases above ~ 573 K (Figure 5b). Such behavior is correlated with the sudden upturn of the σ. The maximum S (Smax) values range between 515 and 550 μV K-1, which is similar to that of undoped SnSe single crystals.31 The power factors (PFs) of Sn1-xPbxSe remain almost constant up to ~ 523 K and suddenly rise afterwards (Figure 5c). The temperature-dependent total thermal conductivity (κtot) of Sn1-xPbxSe gradually decreases with rising temperature, followed by a slight increase afterwards because of bipolar conduction (Figure 5d) as the σ. Importantly, the upturn temperature of the κtot gradually diminishes with higher x: 723, 623, 573, and 523 K for samples with x = 0.05, 0.1, 0.15, and 0.2, respectively. This behavior is strongly related with the reduced Tc driven by PbSe alloying. Figure 5e demonstrates that the κtot is dominated by phonon contribution for all samples. Those with higher x show the lower κlatt because of better scattering thermal phonons by point defects at the atomic level. Nonetheless, the κtot value of Sn1-xPbxSe outpaces that of pristine SnSe above the upturn temperature because the bipolar conduction of the former initiates at lower temperature and raises its κtot higher than that of the latter. Hence this adverse effect overwhelms contribution from phonon scattering at high

Optimization of thermoelectric properties We introduced sodium as a hole-dopant to the Sn1-xPbxSe system to compensate for deficient hole density (nH) caused by PbSe alloying. Doping 1% Na to Sn1-xPbxSe forms Na0.01(Sn1-xPbx)0.99Se phases to exhibit markedly enhanced nH values at 300 K, approximately two orders of magnitude larger than that of Sn1-xPbxSe, for example, 4.31  1019 cm-3 for Na0.01(Sn0.95Pb0.05)0.99Se (Figure S9a). The nH lessons gradually with increasing x as observed in Sn1-xPbxSe. 6



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Figure 6. Temperature dependence of (a) electrical conductivity, (b) Seebeck coefficient, (c) power factor, (d) total thermal conductivity, (e) lattice thermal conductivity, and (f) figure of merit ZT of Na0.01(Sn1-xPbx)0.99Se (x = 0 – 0.2) Despite boosting the nH, Na doping seriously deteriorates the μH: ~ 6 cm2 V-1 s-1 for the sample with x = 0.05 in comparison with 28 cm2 V-1 s-1 of undoped SnSe (Figure S9b). Similar behavior is also observed when alkali-ion is doped to polycrystalline pristine SnSe.40 This result implies that improving the μH would be a next task for further advance in heavilydoped polycrystalline SnSe thermoelectrics.

The temperature of the Smax shifts downward with increasing x, consistent with a trend of the σ. To understand the effect of PbSe alloying on the S, we calculated the Pisarenko relation between the S and the nH for SnSe, and compared the results with the experimental S values at 300 K for Sn0.96Pb0.04Se and Na0.01(Sn0.95Pb0.05)0.99Se in this work as well as for the previous report.31,37,40 (See supporting information and Figure S10 for details) The results demonstrate that PbSe alloying marginally affects S values, which is consistent with the results of our electronic structure calculations showing that it negligibly disturbs the VBM.

The enhanced nH remarkably improves the σ to 59 S cm-1 for Na0.01(Sn0.95Pb0.05)0.99Se from 0.8 S cm-1 for Sn0.95Pb0.05Se at 300 K (Figure 6a). The σ increases from 300 to 373 K, subsequently decreases up to 673 K, and rises quickly afterwards due to bipolar conduction. Very importantly, the σ upturn in the third region is driven by the Pnma-Cmcm transition as discussed in the previous section. Shifting the Tc downward by PbSe alloying reduces the onset temperature of bipolar conduction and resultantly the upturn temperature of the σ value. Earlier activation of bipolar conduction contributes to the enhanced σ and PF at operating temperature of TE conversion of SnSe-based materials and accordingly improved average ZT. For example, the σ upturns at 673, 623, 623, and 573 K for samples with x = 0.05, 0.1, 0.15, and 0.2, respectively. This effect is not obviously observed for Sn1-xPbxSe due to their too low σ values.

The maximum PF value for Na0.01(Sn0.99Pb0.05)0.99Se is 7.39 μW cm-1 K-2 at 823 K, which is remarkably improved from that of Sn0.95Pb0.05Se and Na0.01Sn0.99Se, mainly due to the gain in the σ value (Figure 6c). The κtot and the κlatt of the Na0.01(Sn1-xPbx)0.99Se show the similar trend to those of Sn1-xPbxSe with the comparable value at 300 K (Figure 6d). The minimum κtot of the former is lower than that of the latter: ~ 0.44 W m-1 K-1 at 773 K of Na0.01(Sn0.95Pb0.05)0.99Se compared to ~ 0.60 W m-1 K1 at 773 K of Sn0.95Pb0.05Se. The degree of PbSe alloying fraction influences on a temperature dependence of the κlatt. The κlatt of samples with x = 0, 0.05, and 0.1 follows a power series of T-1.2, T-1.1, and T-0.7, respectively, revealing that the Umklapp phonon scattering is dominant for x < 0.1 (Figure 6e). The sample with x = 0.05 exhibits the highest ZT of 1.16 at 773 K among the series, which is attributed

The S of Na0.01(Sn1-xPbx)0.99Se increases with temperature from RT to 600 – 673 K and subsequently decreases afterwards (Figure 6b). The sample with x = 0.05 exhibits the largest Smax of 378 μV K-1 at 673 K. 7


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(T) for all samples is temperature-dependent. Above ~ 475 K, the RH (T) of Na-doped samples of Na0.01Sn0.99Se and Na0.01(Sn0.96Pb0.04)0.99Se increases gradually while that of the others declines rapidly. For undoped samples, the carrier concentration (nH) is the order of ~ 1017 cm3 at RT and slowly increases until 523 K, followed by rapid increase afterwards because of the intrinsic carrier excitation (Figure 8b). This trend is related with the rapid increment of the σ above 523 K. The nH for Na-doped samples is the order of ~ 1019 at RT with a decrease above 523 K. This is due to scattering between major carriers of high concentration.40 Pb alloying lowers the nH because Pb substitution for Sn reduces intrinsic Sn vacancies that are the major source of charge carriers in p-type SnSe.

simultaneously from the remarkably enhanced σ and the lowered κtot with the least loss of the S compared to Sn1-xPbxSe. (Figure 6f). After obtaining the high ZT of 1.16 from the sample of x = 0.05, we finely tuned the Pb level in the range of x = 0.03 0.06 for optimizing TE performance. The samples generally show a similar trend for TE properties versus temperature (Figure S11). The x = 0.04 sample displays the highest PF of 6.7 μW cm-1 K-2 and the lowest κtot of ~ 0.43 W m-1 K1 , resulting in the highest ZT of 1.18 among this family of Na0.01(Sn1-xPbx)0.99Se, slightly larger than 1.16 of the x = 0.05 sample (Figure 7). Despite the polycrystalline nature of the samples, their TE properties are highly anisotropic. The ZT taken parallel to the press direction of SPS is lower at 0.88 (Figure S12). The achieved ZT is improved by a twofold magnitude from ZT = 0.6 of pristine SnSe polycrystals40 as well as by 43 % from ZT = 0.8 of Na0.01Sn0.99Se40 because of the simultaneous contributions from enhanced PF due to tuning the Tc and the reduced κlatt due to effective point defect scattering by PbSe alloying. Importantly, PbSe alloyed and Na-doped SnSe in this work shows remarkably enhanced average ZT from RT to 773 K: ZTave ~ 0.5 for the sample with x = 0.05 in comparison with ~ 0.35 for Na0.01Sn0.99Se40 and phase-separated Sn1-xPbxSe.21 Recently, ball-mill-processed K-doped SnSe was reported to show the ZT of 1.1 and ZTave ~ 0.5.41 However, we could not enhance ZT of the Sn1xPbxSe system by doping either K or Na and K together.

Figure 8c represents the hole mobility (H) as a function of temperature. For all samples, it increases from RT to 373 K. This trend is attributed to the barrier-like scattering due to defects in grain boundaries as reported earlier.40,62 Afterwards, the H for all samples follows a temperature dependence of T1.5 from 373 K to 523 K. This is the characteristic behavior given by acoustic phonon scatterings.62,63 The H of Na-doped samples continuously decreases above 523 K because scattering between carriers of high concentration becomes significant at high temperatures.64 In contrast, the H of undoped samples increases with rising temperature, attributed to excitation of carriers. Note that PbSe alloying moderately suppresses the H because of the point defect scattering.46 Scanning (STEM)

Transmission

Electron

Microscopy

The AgPbmSbTe2+m (LAST-m) system had long been regarded as a solid solution between isostructural PbTe and AgSbTe2.65 However, its members with m > 10 were revealed to be a bulk material that spontaneously forms a considerable amount of nanostructures even though they appear single phase by conventional XRD, following Vegard’s law. The LAST-m is the first nanostructured bulk thermoelectric material reported, opening a new paradigm of thermoelectrics, of which nanostructuring is a key aspect of high TE performance due to substantial reduction in κlatt.18,66,67 This example underscores the importance of atomic resolution investigations for properly understanding thermal and electrical transport properties of TE materials.

Figure 7. Temperature-dependent figure of merit ZT of Na0.01(Sn1-xPbx)0.99Se (x = 0, 0.03 – 0.06). Temperature-dependent Hall effect measurements We performed the temperature-dependent Hall effect measurements for SnSe, Na0.01Sn0.99Se, Sn0.96Pb0.04Se, and Na0.01(Sn0.96Pb0.04)0.99Se to better understand the effect of Pb and Na incorporation on electrical transport properties. The Hall coefficient RH

We investigated detailed nanostructures for the SPS Na0.01(Sn0.96Pb0.04)0.99Se sample employing atomic resolution STEM, EDS, and EELS. Typical cross-sectional bright-field (BF) STEM image viewed 8



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Figure 8. (a) Hall coefficient, (b) carrier concentration and (c) mobility of SnSe, Sn0.96Pb0.04Se, Na0.01Sn0.99Se, and Na0.01(Sn0.96Pb0.04)0.99Se as a function of temperature. consistent with the result in Figure 9c.

down the [010] axis demonstrates the presence of nanostructures with a size of ~ 5 10 nm, evenly distributed in the matrix. The Na0.01(Sn0.96Pb0.04)0.99Se sample (Figure 9a) contains more distinct and denser nanostructures than the Sn0.96Pb0.04Se (Figure S13). The selected area electron diffraction (SAED) pattern (inset, Figure 9a) on the former is clearly indexed as the SnSe structure along the [010] zone axis, indicating the nanostructures and the matrix are isostructural. Absence of extra diffraction spots excludes the presence of second-phase precipitates such as PbSe and its structural analogues. Highmagnification annular BF STEM (ABF-STEM) image focusing on the individual nanostructure reveals a nearly coherent interface with the matrix (Figure 9b). The corresponding fast Fourier transform (FFT) image does not show extra spots deviating from the SnSe structure (inset, Figure 9b). The inverse FFT (IFFT) image of the (101) atomic planes clearly shows edge dislocations marked by the red symbols in Figure 9c.

The cross-sectional HAADF-STEM image taken on the nanostructure viewed down the b-axis clearly shows layered crystal structure of Na0.01(Sn0.96Pb0.04)0.99Se (Figure 9e). Because the contrast of HAADF-STEM image is approximately proportional to the square of the atomic number (Z),68 apparently bigger and brighter spheres can be assigned to heavier Sn or Pb atoms and weaker to Se atoms. Individual double atomic layers consisting of short (Sn, Pb)–Se bonds are stacked along the a-axis. The slabs are bound by long (Sn, Pb)Se interactions. Figure 9f displays atomic resolution elemental mapping by STEM-EDS for the area marked by a red rectangle shown in Figure 9e. Signals of Se, Sn, and Pb atoms are depicted in red, yellow and blue colors, respectively, in the lower panels in Figure 9f, jointly creating the image in the upper panel. In the latter, Sn and Pb atoms in whitish color are verified to occupy the same atomic sites and be totally disordered while Se atom in red color takes its own positions (Figure 9f). The atomic arrangement generated by elemental mapping given here is consistent with crystal structure of SnSe. Note that segregated PbSe precipitates are not found therein.

Afterwards, we analyzed the compositional difference between the nanostructures and the matrix using the STEM-energy dispersive X-ray spectroscopy (EDS). We observed that Pb is more abundant in the nanostructures, whereas Sn is richer in the matrix with similar abundance of Se in both the regions (Figure S14). Both high-magnification atomic resolution high-angle annular dark-field (HAADF) and ABF-STEM images in Figure 9d support the STEM-EDS results. Examined along the yellow dashed line in Figure 9d, the nanostructure region gives a stronger signal intensity of the Z-contrast in ABF-STEM (vice versa in HADDF-STEM image) than the matrix due to the higher degree of Pb content. A high-magnification ABF-STEM image focusing on the nanostructure shows severely distorted atomic arrangements compared with that on the matrix probably because of compositional discrepancy and edge dislocations around the former (Figure S15),

Because Na is much lighter than other constituent elements, its position could not be determined by the STEM-EDS analysis. Accordingly, we performed a STEM-EELS elemental scan profile at the atomic level to confirm the presence and the relative location of Na atoms in the Na0.01(Sn0.96Pb0.04)0.99Se sample. Figure 9g shows the Na L2,3 edge of the EELS spectrum at 31 eV verifying the presence of Na atoms in the sample, which is absent from that of Sn0.96Pb0.04Se. The line scan profile of the Na EELS spectrum across the nanostructure and the matrix demonstrates that Na concentration is lower in the former (Figure 9h). Overall, combined analysis of STEM, EDS, and 9


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Figure 9. Cross-sectional STEM images and elemental mapping of Na0.01(Sn0.96Pb0.04)0.99Se. (a) Typical lowmagnification bright-field (BF) STEM image. Inset: selected area electron diffraction pattern along the [010] axis. (b) Annular BF(ABF)-STEM image focused on a nanostructure in (a). Inset: fast Fourier transform (FFT) image along [010] axis. (c) Inverse FFT image of (b). Red symbols indicate edge dislocations. (d) High-angle annular dark-field (HAADF) STEM (upper) and ABF-STEM (lower) atomic-resolution images focused on a nanostructure. Insets: Signal intensity of Z-contrast across a nanostructure taken by a line profile. (e) Magnified atomic-resolution HAADF-STEM image taken on a matrix, showing crystal structure of Sn1-xPbxSe at the atomic scale. White arrows indicate spaces between two atom-thick Sn1-xPbxSe slabs. (f) Elemental mapping of Se, Sn, and Pb atoms depicted in red, yellow, and blue color, respectively, by STEM energy dispersive X-ray spectroscopy scanned on the region of a red rectangle in (e) (lower panels), confirming the location of respective atoms in crystal structure of Sn1xPbxSe. These three panels jointly create the image in the upper panel, verifying Sn and Pb atoms are completely disordered in Sn1-xPbxSe. (g) STEM-electron energy loss spectroscopy (EELS) signals for samples of Sn0.96Pb0.04Se and Na0.01(Sn0.96Pb0.04)0.99Se. The black arrow indicates an energy of Na L2,3 edge. (h) EELS signal intensity of Na L-edge across a nanostructure recorded by a line profile. EELS results reveal that PbSe alloying spontaneously induces a considerable amount of nanostructures of Sn1-xPbxSe despite appearance as single phase by conventional XRD. The nanostructures contain relatively more Pb and less Sn and Na concentrations than the matrix. This

observation is in striking contrary to the published phase diagram and literature46,47 that describe Sn1xPbxSe as solid solution. The slightly different chemical composition and indistinguishable unit cell dimension of the nanostructure from the matrix give rise to highly coherent interface between them. Sn and 10



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hydrothermally prepared Sn1-xPbxSe nanoparticles containing PbSe precipitates.45

Pb atoms are totally disordered. As a result, point defects, lattice strains, and mass fluctuation between the nanostructures and the matrix caused by PbSe alloying and Na doping can synergistically interrupt short-wavelength phonon transfer with simultaneously enhanced hole carrier transport, plausibly explaining high TE performance of Na0.01(Sn0.96Pb0.04)0.99Se.65

CONCLUSIONS We demonstrate that PbSe alloying causes phase transition of SnSe to occur at low temperatures, thereby decreasing the onset temperature of bipolar conduction. Interestingly, its effect on a ZT is ambivalent. When carrier concentration is too low as 1017 cm-3), its earlier Sn1-xPbxSe (nH ~ 1016 initiation predominantly contributes to increasing thermal conductivity at elevated temperatures and deteriorates a ZT. On the other hand, raising the nH to the level of ~ 1019 cm-3 by 1% Na doping, dramatically enhances electrical conductivity, giving remarkably improved power factor. Simultaneously, lattice thermal conductivity is reduced by Na doping due to point defect scattering. As a result, the highest ZT of ~ 1.2 among bulk-processed, polycrystalline SnSe thermoelectrics is achieved for Na0.01(Sn0.96Pb0.45)0.99Se, which is two-fold and ~ 40% magnitude larger than that of pristine SnSe and Na0.01Sn0.99Se, respectively. Na0.01(Sn0.95Pb0.05)0.99Se exhibits significantly improved average ZT of ~ 0.5 in the temperature range from RT to 773 K, which is also the highest reported for polycrystalline SnSe thermoelectrics. Importantly, employing atomic resolution scanning transmission electron microscope, we also discover that PbSe alloying spontaneously generates substantial amounts of nanostructures of Sn1-xPbxSe with compositional fluctuation and lattice strains across the interfaces rather than forms solid solutions as long believed, further reducing lattice thermal conductivity. Accordingly, impurity alloying can favorably affect both electrical and thermal transport behaviors and thereby serve as a new platform to enhance a ZT of SnSe thermoelectrics.

Figure 10. The calculated latt with application of various scattering mechanisms in comparison with the experiment values for the samples of pristine SnSe, Pb alloyed Sn0.95Pb0.05Se, Na doped and Pb alloyed Na0.01(Sn0.95Pb0.05)0.99Se. To support this, we performed theoretical calculations on lattice thermal conductivities for the samples of pristine, Pb alloyed, and Na-doped and Pb alloyed SnSe with consideration of various wellknown scattering mechanisms based on the models of Umklapp (U),69,70 Umklapp combined with point defects (U+PD),71,72 U+PD with interfaces, and U+PD with interfaces and dislocations73 (See Supporting Information for further details). Figure 10 shows the calculated latt curves in comparison with experimental values. The calculation results clearly demonstrate that experimental latt values for Na doped and Pb alloyed samples are located well below the latt curve based on the U+PD model over the entire temperature range, indicating multiple phonon scatterings occur. Considering both nanostructuring and dislocations observed in STEM studies in this work as the significant source of phonon scattering, our theoretical model fits well with experimental values. As a consequence, the synergistic effect of nanostructuring and dislocations could be the main reason of the reduced latt observed for Na-doped and Pb-alloyed samples, rather than single point defect scattering. However, reduction in latt is not as dramatic as shown in representative nanostructured bulk thermoelectric materials20,24,66,74,75 and

ASSOCIATED CONTENT Supporting Information Data of elemental analyses, XRD, optical absorption, DFT calculations, Hall effect measurements, TE properties, and additional STEM images. This material is available free of charge via the Internet at http://pubs.acs.org. AUTHOR INFORMATION Corresponding Author *[email protected] Notes The authors declare no competing financial interest. ACKNOWLEDGEMENTS 11


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Figure 1. Powder XRD patterns of (a) Sn1-xPbxSe (x = 0 – 0.3) and (b) Na0.01(Sn1-xPbx)0.99Se (x = 0 – 0.2) materials with standard diffraction patterns of SnSe (The International Centre for Diffraction Data (ICDD) PDF 01-075-6133) and PbSe (00-006-0354). PbSe Bragg peaks are marked by asterisks. 363x465mm (72 x 72 DPI)


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Figure 2. A bulk phase diagram of the PbSe-SnSe ternary system adapted from ref 42. 381x340mm (72 x 72 DPI)



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Figure 3. In situ temperature-dependent PXRD patterns of (a) SnSe, (b) Sn0.95Pb0.05Se, and (c) Sn0.9Pb0.1Se heated from 740 K to 830 K with the interval of 10 K, showing temperature-driven phase transition from the Pnma to Cmcm space group. The (111) reflection peak of the Pnma is gradually weakened on heating. 388x247mm (72 x 72 DPI)


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Figure 4. Differential scanning calorimetry analysis of Sn1-xPbxSe (x = 0 – 0.2) at a rate of 10 K min-1 on heating, showing an endothermic peak for temperature-driven phase transition. 316x270mm (72 x 72 DPI)



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Figure 5. Temperature dependence of (a) electrical conductivity, (b) Seebeck coefficient, (c) power factor, (d) total thermal conductivity, (e) lattice thermal conductivity, and (f) figure of merit ZT of Sn1-xPbxSe (x = 0 – 0.2). 667x387mm (72 x 72 DPI)


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Figure 6. Temperature dependence of (a) electrical conductivity, (b) Seebeck coefficient, (c) power factor, (d) total thermal conductivity, (e) lattice thermal conductivity, and (f) figure of merit ZT of Na0.01(Sn1xPbx)0.99Se (x = 0 – 0.2). 665x386mm (72 x 72 DPI)



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Figure 7. Temperature-dependent figure of merit ZT of Na0.01(Sn1-xPbx)0.99Se (x = 0, 0.03 – 0.06). 293x258mm (72 x 72 DPI)


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Figure 8. (a) Hall coefficient, (b) carrier concentration and (c) mobility of SnSe, Sn0.96Pb0.04Se, Na0.01Sn0.99Se, and Na0.01(Sn0.96Pb0.04)0.99Se as a function of temperature. 634x278mm (72 x 72 DPI)



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Figure 9. Cross-sectional STEM images and elemental mapping of Na0.01(Sn0.96Pb0.04)0.99Se. (a) Typical lowmagnification bright-field (BF) STEM image. Inset: selected area electron diffraction pattern along the [010] axis. (b) Annular BF(ABF)-STEM image focused on a nanostructure in (a). Inset: fast Fourier transform (FFT) image along [010] axis. (c) Inverse FFT image of (b). Red symbols indicate edge dislocations. (d) Highangle annular dark-field (HAADF) STEM (upper) and ABF-STEM (lower) atomic-resolution images focused on a nanostructure. Insets: Signal intensity of Z-contrast across a nanostructure taken by a line profile. (e) Magnified atomic-resolution HAADF-STEM image taken on a matrix, showing crystal structure of Sn1-xPbxSe at the atomic scale. White arrows indicate spaces between two atom-thick Sn1-xPbxSe slabs. (f) Elemental mapping of Se, Sn, and Pb atoms depicted in red, yellow, and blue color, respectively, by STEM energy dispersive X-ray spectroscopy scanned on the region of a red rectangle in (e) (lower panels), confirming the location of respective atoms in crystal structure of Sn1-xPbxSe. These three panels jointly create the image in the upper panel, verifying Sn and Pb atoms are completely disordered in Sn1-xPbxSe. (g) STEM-electron energy loss spectroscopy (EELS) signals for samples of Sn0.96Pb0.04Se and Na0.01(Sn0.96Pb0.04)0.99Se. The black arrow indicates an energy of Na L2,3 edge. (h) EELS signal intensity of Na L-edge across a nanostructure recorded by a line profile. 1116x1023mm (72 x 72 DPI)


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Figure 10. the calculated κlatt with application of various scattering mechanisms in comparison with the experiment values for the samples of pristine SnSe, Pb alloyed Sn0.95Pb0.05Se, Na doped and Pb alloyed Na0.01(Sn0.95Pb0.05)0.99Se 310x245mm (72 x 72 DPI)



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TOC 719x283mm (72 x 72 DPI)


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Enhancing p-type thermoelectric performances of polycrystalline SnSe

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