Polycrystalline SnSe–Sn1–vS solid solutions: Vacancy engineering and nanostructuring leading to high thermoelectric performance

Journal Pre-proof Polycrystalline SnSe–Sn1–vS Solid Solutions: Vacancy Engineering and Nanostructuring Leading to High Thermoelectric Performance Asfandiyar, Bowen Cai, Hua–Lu Zhuang, Huaichao Tang, Jing–Feng Li PII:

S2211-2855(19)31107-3

DOI:

https://doi.org/10.1016/j.nanoen.2019.104393

Reference:

NANOEN 104393

To appear in:

Nano Energy

Received Date: 17 September 2019 Revised Date:

25 November 2019

Accepted Date: 8 December 2019

Please cite this article as: Asfandiyar, B. Cai, H.–L. Zhuang, H. Tang, J.–F. Li, Polycrystalline SnSe– Sn1–vS Solid Solutions: Vacancy Engineering and Nanostructuring Leading to High Thermoelectric Performance, Nano Energy, https://doi.org/10.1016/j.nanoen.2019.104393. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2019 Elsevier Ltd. All rights reserved.

Graphical abstract Sn vacancies and further Ag doping at the sites of Sn are introduced into the SnSe–Sn1-vS matrix by a facile technique, leading to significant power factor (PF) improvement (5.3 µW cm–1 K–2) and the simultaneous reduction in thermal conductivity (≈0.250 Wm−1 K−1). As a result, high ZT ≈1.75 at 823 K was achieved in Sn0.978Ag0.007S0.25Se0.75 sample.

Polycrystalline SnSe–Sn1–vS Solid Solutions: Vacancy Engineering and Nanostructuring Leading to High Thermoelectric Performance Asfandiyar, a Bowen Cai, a Hua–Lu Zhuang, a Huaichao Tang, a and Jing–Feng Li a* a

State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and

Engineering, Tsinghua University, Beijing 100084, China.

Abstract P–type polycrystalline Sn1-vS compounds and Sn0.985S1-xSex solid solutions are prepared by combining mechanical alloying (MA) and spark plasma sintering (SPS). Dislocations in the grains are successfully introduced through vacancy engineering in Sn1-vS compounds. Point defects provide high–frequency phonons scattering from the Se substitution in Sn0.985S1-xSex, in addition to the strong scattering of mid– frequency phonons by dislocations. This leads to a low lattice thermal conductivity and an enhanced thermoelectric figure of merit (ZT) of ≈1.1 for the Sn0.985S0.25Se0.75 at 823 K. The Ag dopant is selected to further enhance the electrical transport properties of the optimized composition. The power factor improved from 4.5 (Sn0.985S0.25Se0.75) to ≈5.3 µW m−1 K−2 for Sn0.978Ag0.007S0.25Se0.75 sample. Surprisingly, the Ag doping induced a nanostructured matrix with dispersed spherical coherent precipitates of AgSnSe2 inside the grains, which further strengthens the scattering of phonons. The presence of AgSnSe2 nanoscale precipitates inside the grains and dislocations–induced by Ag at the grain boundaries contributes to an impressively low lattice thermal conductivity (κL). Consequently, the maximum ZT ≈1.75 at 823 K is achieved for the Sn0.978Ag0.007S0.25Se0.75 sample. Keywords SnSe-Sn1-vS; Mechanical alloying; Spark plasma sintering; Thermoelectric properties Correspondence to: [[email protected]]

1. INTRODUCTION The fossil fuels shortage and increasing related environmental problems (emissions of any greenhouse gas) are red–alarm for the world’s increasing energy demands. Thermoelectricity, a physical phenomenon that converts heat directly to electricity and vice versa, provides a potential solution to the world energy crisis as well as to environmental issues that arise from fossil fuels [1]. The thermoelectric conversion efficiency is generally measured by the temperature dependent dimensionless figure of merit ZT = S2σT/ κT = S2σT/(κe + κL) [2], where σ, S, S2σ, T, κT, κe and κL are the electrical conductivity, Seebeck coefficient, power factor, absolute temperature, total thermal conductivity, the electronic and lattice thermal conductivity, respectively [3-5]. A good thermoelectric material requires a high numerator (S2σ) and low denominator (κT) values in the above equation. It is hard to enhance the ZT values owing to the strong coupling between S, σ and κe as S inversely, and σ and κe are directly proportional to the carrier concentration. To date, two major strategies for achieving high ZTs are to promote the electrical performance via increasing the power factor (S2σ) and to reduce the independent material property, lattice thermal conductivity (κL) [6]. The former strategy has been demonstrated by the band engineering concept [7,8], charge carrier engineering [9-11] and chemical doping [12-14] while the later one can be accomplished through introducing phonons scattering by a few 1

approaches that include Nanostructuring [15,16], mesoscale grains and boundaries [17], artificial induced nanoscale dislocations [18], atomic–scale point defects [19] and multiscale hierarchical architectures [20]. Lead free, binary tin chalcogenides family (SnX; X = Se, S and Te) have received increasing attention as a common available, hazardless and cheap thermoelectric materials particularly SnSe and SnS [21]. Among the various promising thermoelectric materials, stannous selenides (SnSe) has attracted strong interest since few years ago, because a peak ZTs of ≈2.8 at 773 K and ≈2.6 at 923 K were reported in its pristine n–type [22] and p–type [23] single crystalline form, owing to their high power factor and remarkably low thermal conductivity values along specified crystal directions. Moreover, a broad ZT plateau over ≈1 along the b– crystallographic direction was achieved in hole–doped SnSe single crystals [24]. In general, single crystals have higher thermal conductivity than polycrystals [25,26]. Besides, SnSe in its single crystalline form are not well–suited for thermoelectric devices due to the high cost for production, weak cleaving properties and special demands of crystal–growth technique [27,28]. Therefore, polycrystalline SnSe [29-31] has become a promising alternative candidate for practical applications due to machinability and scalability. The highest ZTs of pristine polycrystalline SnSe verses its single crystalline counterparts are much lower because polycrystalline samples have strong anisotropic nature, and possess low carrier mobility and higher thermal conductivity, which stem from the effects of Sn oxides [32-34]. To improve the electrical conductivity of polycrystalline SnSe, several major strategies have been employed, such as Sn vacanciesc [35], dual (Ag/Na) [36], K [37], alkali metal (Li, Na, K) [38,39], Ag and Na doping at Sn sites [40]. Few effective strategies have been adopted in reducing the independent material property (κL) of polycrystalline SnSe, even though secondary phases such as Ag8SnSe6 [36], AgSnSe2 [41], PbSe [42], SnTe [43] and ZnSe [44], have been used. Stannous sulfide (SnS), with an analogous structure to SnSe, has been recognized as a potentially high– performance thermoelectric material [45]. SnS in its single [46] and polycrystalline [47] form performs rather poorly as a result of its low power factor and high thermal conductivity than that of the SnSe [23, 24, 36]. Therefore, in the past very few strategies including Ag [48] and Na [49] doping have been employed to upsurge its electrical conductivity. Similarly, very less effective strategies, such as SnS–SnSe [50] solid solutions and further doping with Ag at Sn sites have been proposed to reduce the thermal conductivity. Up to date, there is no effective strategy used in reducing the independent lattice thermal conductivity of polycrystalline SnS and their solid solutions (SnS)1-x–(SnSe)x. Besides, no sufficient discussions were made on its microstructural observations, which could fully, support the experimental data for the already reduced thermal conductivity. Moreover, ZT has been improved to ≈1.3 for polycrystalline Na [51] doped SnSexS1-x before 850 K. However, samples doped with Na are not stable upon repeated heating and cooling [41]. Furthermore, SnSe become unstable due to the volatilizing of Sn and/or Se atoms at elevated temperature, so it has been a challenge to enhance ZT even higher to ZT >1.5 before 850 K, which is essential for device applications. In this work, we synthesized Sn1-vS compounds with different nominal Sn off stoichiometric ratios, pure SnSe and solid solutions among the optimized (Sn0.985S) 1–x–(SnSe) x compositions. The spark plasma sintered (SPS–ed) Sn1-vS samples are found to exhibit dislocations within the grains, leading to significantly enhanced mid–frequency phonons scattering. In the solid solutions S and Se induced mass and strains fluctuations also favors the scattering of high–frequency phonons. Furthermore, Ag was chosen as a p–type dopant to improve the electrical transport properties of optimized Sn0.985S0.25Se0.75 sample, and the power factor was lifted up significantly from 4.5 (Sn0.985S0.25Se0.75) to ≈5.3 µW m−1 K−2 for Sn0.978 Ag0.007S0.25Se0.75 sample. In addition, the structural and morphological characterizations including X–ray diffraction (XRD), scanning transmission electron microscopy (STEM) with energy dispersive spectroscopy (STEM–EDS) and atomic resolution high angle annular dark field–scanning transmission electron microscopy (HAADF– 2

STEM) indicate that Ag doping can result in single crystalline spherical coherent nanoscale precipitates of AgSnSe2 in the matrix and dislocations at grain boundaries, respectively. The AgSnSe2 exhibits an important effect on lattice thermal conductivity and further leads to an extra low κL and in turn a high peak ZT ≈1.75 at 823 K, which is a record value for Sn0.978 Ag0.007S0.25 Se0.75 sample at this temperature compared with previously reported studies [50,51]. To the best of our knowledge, this is the first attempt by using vacancy engineering in pristine SnS and successful nanoprecipitations in their solid solutions that achieves a high record ZT so far in the ball–milled and SPS–ed sample.

2. RESULTS AND DISCUSSIONS 2.1 Phase Figure 1a shows the X–ray diffraction (XRD) patterns of the SnS, Sn0.99S, Sn0.985S, Sn0.98S, Sn0.985S0.5Se0.5 and Sn0.985S0.25Se0.75 SPS–ed bulk samples. The main peaks in the XRD patterns for Sn1-v S samples can be well indexed to orthorhombic crystal structure. A small horizontal shift in 2θ values of diffraction peak (111) toward the lower angle was observed in Sn0.985S1–xSex solid solutions with increasing x, which can be attributed to the smaller ionic radii of S (1.84 Å) than Se (1.91 Å). The lattice parameters (a, b and c) of SnS are: a = 11.201 Å, b = 3.988 Å, and c = 4.332 Å and show negligible change for Sn0.99S, Sn0.985S and Sn0.98S. The lattice parameters of Sn0.985S1–xSex samples were calculated from the XRD patterns, and were plotted in Figure S1a (supporting information). It can be seen that the lattice parameters of Sn0.985S0.5Se0.5 and Sn0.985S0.25Se0.75 samples expand linearly with increasing Se contents (x), following the Vegard’s law, which confirmed the fact that Se take S atoms positions. The XRD patterns taken from SPS–ed Sn0.985dAgdS0.25Se0.75 (d=0.004. 0.007, 0.01) bulk samples are shown in Figure S1b (supporting information). The main peaks in the XRD patterns for Sn0.981Ag0.004S0.25Se0.75, Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 samples can be well indexed to orthorhombic crystal structure. No shift in 2θ was observed of all peaks for Ag–doped vacancy–containing samples in comparison to their undoped counterpart Sn0.985S0.25Se0.75 and the lattice parameters remain unchanged due to the similar ionic radii between Ag+ (1.15 Å) and Sn2+ (1.12 Å). In addition, detectable peaks denoted by blue arrows belonging to the cubic AgSnSe2 structure (JCPDS #331194) [41] were observed in the expanded XRD pattern of the Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 samples as shown in Figure 1b. 2.2 Electronic transport properties Figure 2a shows the temperature dependence of electrical conductivity (σ) for Sn1−vS, pure SnSe and Sn0.985S1–xSex (x = 0.5 and 0.75) SPS–ed samples along the direction perpendicular to SPS pressurizing direction. In the whole temperature range, σ for all the samples shows the same temperature–dependence semiconducting transport behavior. σ of the Sn1-vS samples increases with rising temperature. At room temperature (≈300 K) the σ of pristine SnS is comparable to our previously published data [47] along the specified direction, and increases for Sn0.99S and Sn0.985S and then decreases with further increase in vacancy (v). In addition to remarkably enhanced mid–frequency phonons scattering (section; 2.5), the Sn off stoichiometry significantly enhanced σ for Sn1-vS with increase in v compared to that of stoichiometric SnS. To further investigate the electrical transport behavior in Sn1−vS samples, room–temperature Hall measurements are performed. Compared to the stoichiometric SnS, carrier concentration (nH) is greatly enhanced by the intentional introduction of Sn vacancies. The room–temperature nH for the pristine SnS, and vacancy–containing Sn0.99S, Sn0.985S and Sn0.98S samples are ≈0.00134, 0.293, 0.344 and 0.418 × 1018 cm−3, respectively. We attribute this increase in nH to the Sn vacancies, which act as acceptors and in turn increase the hole carrier concentration. Room–temperature Hall carrier concentration (nH) and Hall mobility (µH) for all Sn1−vS samples are presented in Table 1. µH, for the SnS, Sn0.99S, Sn0.985S, and Sn0.98S samples are ≈4.12, 3

2.47, 2.38 and 1.24 cm2 V–1 S–1. Stoichiometric SnS exhibited high µH than that of the vacancy–containing samples. The lower µH of vacancy–containing samples imply that dislocations distress the hole transport within these samples. Therefore, the observed increase in σ of the vacancy–containing samples is attributed to the increase in nH. σ of the Sn0.985S1–xSex at ≈300 K is further increased with increase in Se contents (x), which is mainly attributed to the increase of carrier concentration (Table 1). σ of the pure SnSe is higher than those of the Sn1-vS and Sn0.985S0.5Se0.5 sample due to high nH. It is seen that over the whole temperature range, Sn0.985S0.25Se0.75 show a higher σ value than those of the Sn1−vS, Sn0.985S0.5Se0.5 and SnSe samples. The maximum σ value of ≈37.045 S cm–1 is obtained at 823 K for the Sn0.985S0.25Se0.75 composition. We doped Ag atoms in the optimized Sn0.985S0.25Se0.75 composition to further improve the thermoelectric performance. Figure 2b displays σ as a function of temperature (T) for the undoped Sn0.985S0.25Se0.75 (as a reference from Figure 2a) and Ag–doped Sn0.985-dAgdS0.25Se0.75 (d = 0.004, 0.007 and 0.01) samples. At room–temperature σ increases up on Ag doping. The room temperature σ values are ≈0.227, 0.795, 0.46, and 0.3 S cm−1 for Sn0.985S0.25Se0.75, Sn0.981Ag0.004S0.25Se0.75, Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 samples, respectively. The room–temperature nH and µH for the corresponding samples are ≈1.042, 0.714, 0.588 and 0.438 × 1018 cm−3, and ≈0.386, 4.24, 2.65 and 1.83 cm 2 V–1 S–1 , respectively as shown in Table 2. The σ is abruptly enhanced from ≈0.227 (Sn0.985S0.25Se0.75) to 0.785 S cm–1 (Sn0.981Ag0.004S0.25Se0.75) at ≈300 K, in particular due to the higher µH. The σ value of all the Ag–doped vacancy–containing samples is higher than that of its undoped vacancy–containing counterpart, because of the much higher µH values. The σ values show greater difference for the three Ag–doped samples in the temperature range from 300 to ≈648 K and then tends to be constant because the texture effect does not exist at high temperature. The σ value increases rapidly with rising temperature probably due to a rapid increase in thermally activated carrier concentration. Of note, the rapid increase begins above ≈648 K in the Ag–doped Sn0.985S0.25Se0.75 samples. The Seebeck coefficient (S) of the pristine, and intentionally off stoichiometric Sn1−vS, pure SnSe and Sn0.985S1–xSex samples as a function of temperature is shown in Figure 2c. Consistent with the increase in nH in the vacancy–containing samples, the S values are much lower than that of the pristine SnS at ≈300 K. Such as for the Sn0.98S, which exhibits the lowest S due to its higher nH value. In the whole range of temperature all the vacancy–containing samples show lower S values than that of the pristine SnS and are in consistence to the electrical conductivity values. S of the Sn0.985S1–xSex samples further decrease with increase in Se contents (x), which is due to the further increase in carrier concentration and are consistence with σ values. S of the SnSe is lower than those of the Sn1-vS and Sn0.985S0.5Se0.5 samples due to high carrier concentration. The room–temperature Seebeck coefficients are ≈526.31, 355.18, 346.22, 323.65, 300.58, 315.36 and 258.72 µV K–1 for the SnS, Sn0.99S, Sn0.985S, Sn0.98S, SnSe, Sn0.985S0.5Se0.5 and Sn0.985S0.25Se0.75 samples, respectively. S of the Sn0.99S, Sn0.985S, Sn0.98S, SnSe, Sn0.985S0.5Se0.5 and Sn0.985S0.25Se0.75 samples increases in the temperature range ≈300–698 K and then decreases faster due to a rapid increase in carrier concentration caused by the intrinsic excitation above ≈698 K. The bipolar temperature of pristine SnS is lower than those of the vacancy–containing samples because of the low carrier concentration. Figure 2d displays the temperature dependence of the Seebeck coefficient for undoped and three Ag–doped Sn0.985S0.25Se0.75 samples. The trajectory of S for all the samples suggests that the major charge carriers are hole. At ≈300 K, S is ≈335.76 µV K–1 for the Sn0.981Ag0.004S0.25Se0.75 sample and then increases for Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 samples because of the decrease in carrier concentration. At room temperature S of the Ag–doped vacancy–containing samples shows higher values than that of its undoped vacancy–containing counterpart. As shown in Table 2, the hole concentration decreases by incorporating Ag into the lattice. As the sample Sn0.985S0.25Se0.75 is a p–type semiconductor, which contains Sn vacancies. These intentional vacancies generate empty electronic states and behave like p–type dopants. Up on doping, Ag atoms fill the Sn vacancies and as a result hole concentration decreases in doped samples. Here the intentional vacancies are greater than that of the doping concentratin (v>d). Thus, at low doping 4

levels (Ag=0.004, 0.007 and 0.01), the p–type charge concentration decreases and as a result S increase. For example, the sample Sn0.975Ag0.01S0.25Se0.75 keep lowest hole concentration and show higher S value. The high S increases with increasing temperature around ≈648 K and then abruptly decreases due to the activation of minority carriers (bipolar effect) caused by the temperature dependent reduction of the band gap. Figure 2e shows the power factor (PF = S2σ) values of the Sn1−vS, SnSe and Sn0.985S1–xSex samples versus vacancy (v) and selenium content (x), respectively. S2σ of SnS has been enhanced due to the improved electrical conductivity value in vacancy–containing samples. S2σ for the SnS, Sn0.99S, Sn0.985S, and Sn0.98S samples are ≈1.12, 2.57, 3.35 and 2.65 µW cm–1 K–2 at 873 K, respectively. Furthermore, as a consequence of the higher electrical conductivity in Sn0.985S1–xSex solid solutions, the power factor further rose to maximum value. S2σ of the pure SnSe is higher than those of the Sn1-vS and Sn0.985S0.25Se0.75 samples. S2σ for all the samples increases with increase in temperature. The maximum S2σ values of the Sn0.985S and Sn0.985S0.25Se0.75 samples are ≈3.35 and 4.5 µW cm–1 K–2, which is due to the high electrical conductivity value at 873 K and 823 K, respectively. S2σ of the undoped and Ag–doped vacancy–containing Sn0.985S0.25Se0.75 samples in the temperature range of 300–823 K is shown in Figure 2f. S2σ of the Ag–doped samples have been enhanced as a consequence of improved σ and high S at ≈300 K. Of note, at elevated temperatures (≈748–823 K) the higher S2σ values of the three Ag–doped samples than that of its undoped counterpart are solely resulted from the enhanced electrical conductivity. As the S values of all the Ag– doped samples are lower than that of the undoped Sn0.985S0.25Se0.75 in this temperature range. The maximum S2σ values for Sn0.981Ag0.004S0.25Se0.75, Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 samples are ≈5.27, 5.3 and 5.22 µW cm−1 K−2 at 823 K, which are higher than that of its undoped counterpart (4.5 µW cm−1 K−2). 2.3 Thermal transport properties Figure 3a shows the variation of total thermal conductivity (κT) for the Sn1−vS, pure SnSe compounds and SnS1–xSex solid solutions as a function of temperature along the direction perpendicular to SPS pressurizing direction. A general trend found for SnS and SnSe thermoelectric material is that, κT decreases with increasing temperature. κT of the pristine SnS sample ranges from ≈1.28 to 0.598 Wm−1 K−1 between ≈300 and 873 K. It is important to note that κT of all the vacancy–containing samples are lower than that of the pristine SnS, and the smallest κT is recorded as ≈0.466 Wm−1 K−1 in Sn0.985S at 873 K, 28% of the value for the pristine SnS. κT of the SnSe is lower than those of the Sn1-vS samples and show phase transition above ≈800 K. In the whole temperature range, κT of the undoped solid solutions is lower than those of the Sn1-vS and SnSe samples. It has been shown previously that the electronic part (κe) of κT (κL+ κe) is negligible due to much low value of σ for SnS and the dominant lattice part (κL) of the thermal conductivity accounts for total thermal conductivity [47]. Inset of Figure 3b shows the variation of total thermal conductivity (κT) of Ag– doped Sn0.985S0.25Se0.75 samples in the temperature range ≈300–823 K. κT decreases with increase in Ag doping (d = 0.4 and 0.7%), whereas for smaple Sn0.975Ag0.01S0.25Se0.75 increases with further increase in Ag elements. The minimum thermal conductivity ≈0.250 Wm−1 K−1 at 823 K was achieved for the Sn0.978Ag0.007S0.25Se0.75 sample. The temperature dependent lattice thermal conductivities (κL) of the undoped and Ag–doped vacancy–containing samples were calculated according to κL = κT − κe = κT – LσT, where L is the Lorentz, σ is the electrical conductivity and T is the absolute temperature. We calculated κL for the undoped vacancy–containing Sn0.985S0.25Se0.75 sample, so that to investigate the influence of phonons scattering from the nanoscale precipitates individually in the Ag–doped samples and will be discuss later. κL of doped vacancy–containing samples is suppressed as compared to that of its undoped counterpart as shown in Figure 3b. Of note, in the whole temperature range κL of the Sn0.981Ag0.004S0.25Se0.75 and Sn0.978Ag0.007S0.25Se0.75 is lower than that of its undoped counterpart. We attribute the suppression of κL to the enhanced low–to–middle frequency phonons scattering (section; 2.5). κL for Sn0.981Ag0.004S0.25Se0.75, 5

Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 samples are 36%, 44% and 25% lower than that of the undoped Sn0.985S0.25Se0.75 sample at 823 K. 2.4 Thermoelectric performance Figure 4 and Figure S2 (supporting information) show the ZT values of the Sn1−vS, pure SnSe, undoped vacancy–containing Sn0.985S1–xSex and Ag–doped vacancy–containing Sn0.985S0.25Se0.75 samples. In Sn1−vS samples, dislocations strengthen mid–frequency phonons scattering (section; 2.5), which significantly reduce lattice thermal conductivity and simultaneously enhance electrical conductivity due to increase in carrier concentration and ultimately improved the power factor. A maximum ZT of ≈0.62 is achieved at 873 K along the direction perpendicular to the SPS pressure in Sn0.985S sample, which is ≈293% higher than that of the pure SnS (in the present work and previously published data) [47] and 44% higher than that of the pure SnSe sample. The ZT values were further increased with increasing Se content (x) and a maximum ZT ≈1.1 at 823 K for the composition Sn0.985S0.25Se0.75 was obtained. Herein, we attribute the high ZT to the low lattice thermal conductivity due to additional high–frequency to mid–frequency phonons scattering and the high power factor (4.5 µW cm–1 K,–2) in Sn0.985S0.25Se0.75. The ZT was further significantly enhanced to maximum values ≈1.75, 1.63 and 1.46 at 823 K for the Sn0.978Ag0.007S0.25Se0.75, Sn0.981Ag0.004S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 samples, among which ZT ≈1.75 is one of the highest value for polycrystalline SnS– SnSe based materials reported at this temperature so far by using ball–milling and SPS technique. We attribute the marked enhancement of the thermoelectric figure of merit to the enhanced power factor enabled by Ag doping, and the simultaneous reduction in lattice thermal conductivity by introducing AgSnSe2 nanoscale coherent precipitates and dislocations (section; 2.5). 2.5 Microstructural Characterizations To study microstructures and determine the detailed crystalline information of Sn1-vS and Sn0.985S1–xSex, atomic resolution high angle annular dark field–scanning transmission electron microscopy (HAADF– STEM) and selected area electron diffraction (SAED) analyses were used. Figure 5a is HAADF-STEM image, which reveal nanoscale dislocations with dark contrast inside the grain. Figure 5b is the zoom–in image of the white framed area in Figure 5a, in which dislocations can be witnessed. Dislocation engineering is an effective and practical strategy to suppress the propagation of phonons at intermediate frequency [52]. Based on the detailed HAADF–STEM analyses, the strengthened phonon scatterings can be found to originate from high densities of dislocations in Sn0.985S. Therefore, the resulted lower values of κT for all the vacancy–containing samples can be attributed to significantly reduced κL, suggesting that mid– frequency phonons scattering have been remarkably enhanced. To determine the detailed crystalline information, selected area electron diffraction (SAED) analyses and atomic resolution high angle annular dark field–scanning transmission electron microscopy (HAADF–STEM) were used. Figure 5c is the SAED pattern of the matrix, which can be indexd to SnS taken along [100] zone axis. (right top inset). Besides, HAADF–STEM image was taken inside a grain to determine the detailed crystalline information of Sn0.985S0.25Se0.75 sample. The framed area (solid black) in Figure 5d, displays local lattice distortions, which should be derived from the alloyed Se with different atomic/ionic size of S. Such local lattice distortions and structural variations enhance the high frequency phonons scattering in a given material. Addition to dislocations (mid–frequency phonons scattering), atomic mass variations, atomic size differences between the dopant and substituent and the lattice size of the host structure in solid solutions are the factors, which help intensify high–frequency phonons scattering. Therefore, the further reduction in lattice thermal conductivity is due to the point defects (lattice distortions) caused by different atomic masses of Se (79.86 g/mol) and S (32.07 g/mol), and strain field fluctuations (structural variations) caused by slightly differences between ionic radii (1.91 Å of Se and 1.84 Å of S). Furthermore, the microstructures and compositional 6

analyses for Sn0.978Ag0.007S0.25Se0.75 sample are investigated by using scanning transmission electron microscopy (STEM) and scanning transmission electron microscopy–energy dispersive X–ray spectroscopy (STEM–EDS). Figure 6a is the STEM image for Sn0.978Ag0.007S0.25Se0.75 sample. An abundant amount of spherical shape nanoscale precipitates with bright–contrast are present for the corresponding sample, in which many precipitates with an average size of ≈10–15 nm can be found embedded inside the grains. The brighter contrast of the nanoscale precipitates than matrix signifying the heavier elements chemical component and will be discuss later. The elemental distribution of the framed area (dashed orange) in Figure 6a, examined by STEM–EDS elemental mapping, reveals that element Ag, Sn, S, and Se distribute uniformly all over the matrix as shown in Figure 6b. Ag doping in the form of Sn substitutions increases the defects (lattice distortions and strain field fluctuation caused by different atomic masses and ionic radii between Ag and Sn atoms) in the Sn0.978Ag0.007S0.25Se0.75 solid solution. A typical HAADF–STEM image of the grain inside a small–size spherical nanoscale precipitate is shown in Figure 6c. The formation of spherical nanoprecipitates inside the grains shows that the interfacial energy is dominant in determining the shape of precipitates. The precipitate has a coherent feature i.e., without structure defects, e.g., lattice distortions, dislocations and stacking faults between the matrix and precipitate. Of note, the spherical coherent precipitates exhibit single crystalline nature. EDS analyses of the matrix (EDSM) and precipitate (EDSP) of the framed areas (dashed white and solid yellow) in Figure 6c are shown in Figure 6d and e. According to the EDSM and EDSP, Se intensities are higher than that of the Sn element, which confirm that Sn vacancies were successfully engineered. Moreover, the EDSP shows lower intensity for Sn elements compare with that of the EDSM, which exhibits that these precipitates are rich in Ag (fewer amounts of Sn atoms) and show brighter contrast than the matrix. Besides, from the EDSP analyses, the precipitates consist of Ag, Sn and Se without any S atoms, indicating the composition of precipitate is AgSnSe2. The STEM– EDS analysis also confirms that the precipitates are AgSnSe2, which is consistent with the XRD results. Addition to AgSnSe2, Ag (nanoparticles) mostly segregates at the grain boundaries as shown in the right top inset of Figure 6f and Figure S3 (supporting information). Similar to spherical coherent nanoprecipitates, these nanoparticles shows brighter contrast than matrix indicating richer Ag area. To understand the crystal structure of these nanoparticles, we performed atomic resolution high angle annular dark field–scanning transmission electron microscopy (HAADF–STEM) as shown in Figure 6f. Unlike AgSnSe2, Ag atoms segregate to dislocations at the grain boundaries. The XRD patterns and STEM–EDS elemental mapping and analyses show that Ag elements not only dope in the form of Sn substitutions in the lattice of Sn0.985S0.25Se0.75 sample, but also form nanoscale spherical coherent precipitates of AgSnSe2 in the matrix and some segregate to dislocations at grain boundaries. In the whole temperature range κL of the Sn0.981Ag0.004S0.25Se0.75 and Sn0.978Ag0.007S0.25Se0.75 lower than that of the undoped Sn0.985S0.25Se0.75 sample, which can be attributed partly to the phonons scattering by Ag point defects and the enhanced phonon– electron coupling. Of note, κL for Sn0.981Ag0.004S0.25Se0.75, Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 samples are 36%, 44% and 25% lower than that of the undoped Sn0.985S0.25Se0.75 sample at 823 K, which suggests that the presence of AgSnSe2 coherent nanoscale precipitates plays an important role in the reduction of κL. Our HAADF–STEM analysis suggests that Ag point defects scatter high–frequency phonons (addition to point defects by Se alloying) [53], dislocations induced by Ag scatter mid–frequency phonons (addition to dislocations induced in off stoichiometric Sn0.985S0.25Se0.75) [54,55], while single crystalline coherent nanoprecipitates probably scatter the low–to–middle frequency phonons in Sn0.978Ag0.007S0.25Se0.75 sample [56-58]. 3. Conclusion In summary, we achieved ZT ≈0.62 at 873 K in p–type Sn0.985S, associating with the reduced κL and improved S2σ. High density nanoscale dislocations inside the grains together with atomic–scale point defects 7

realizes the strong mid and high–frequency phonons scattering, accounting for the significantly reduced κL. A higher ZT value of ≈1.1 was obtained at 823 K for the Sn0.985S0.25Se0.75 sample. A nanoscale precipitates and dislocations induced by Ag are effective in suppressing the already low thermal conductivity of the Sn0.985S0.25Se0.75. The maximum 44% of lattice thermal conductivity was suppressed at 823 K. The electrical conductivity was additionally increased due to the increase in carrier mobility up on Ag doping. As a result, further 59% enhanced ZT (1.75) is achieved for Sn0.978Ag0.007S0.25Se0.75. This high performance adds more flexibility to the strategy for enhancing thermoelectric performance by producing nanoscale precipitates and dislocations, which can enhance thermoelectric performance in broader materials. Experimental High purity elements Sn (99.99%), S (99.99%), Se (99.99%) and Ag (99.5%) raw powder were weighed individually in a Argon (Ar) glovebox. Then, the weighted elements in the form of compositions SnS, Sn0.99S, Sn0.985S, Sn0.98S, SnSe, Sn0.985S0.5Se0.5, Sn0.985S0.25Se0.75, Sn0.981Ag0.004S0.25Se0.75, Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 were subjected to the special sealed stainless steel jars containing stainless steel balls as shown in Figure S4 (supporting information). The weight ratio of ball to powder was 20:1. All the compositions were ball–milled at 450 revolutions/minutes in nine jars for 800 minutes. After ball milling all the sealed jars and graphite molds of inner diameter ≈13.5 mm were transferred to the glovebox and filling of the molds were done in Argon atmosphere. Then, all the molds one after another transferred to spark plasma sintering system and were sintered at 850 (SnS, Sn0.99S, Sn0.985S and Sn0.98S) and 810 K (SnSe, Sn0.985S0.5Se0.5, Sn0.985S0.25Se0.75, Sn0.981Ag0.004S0.25Se0.75, Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75) for 10 minutes under an axial pressure of 50 MPa in vacuum. The obtained SPS–ed specimens were cylindrical shape with dimensions of Φ 13 mm × 13 mm. Phase structure was analyzed by XRD (D8 Advance, Bruker, Germany) with Cu–Kα radiation. The microstructure of the bulk sample was confirmed by transmission electron microscopy (TEM, JEOL, JEM 2010F, Japan). Two bulk specimens were obtained via cutting and grinding, one is disk–like and about ≈ 10 mm × 1 mm for thermal conductivity measurements and the second is strip–like about ≈ 2.1 mm × 2.1 mm × 10 mm for Seebeck coefficient / electrical resistivity measurement. SnS and SnSe crstalize in anisotropic crystal structure and generally exihibits high thermoelectric performance along the direction parallel to SPS pressure [47]. However, polycrystalline specimens cut along the direction parallel to SPS pressure and their thermoelectric properties along the specified direction are not stable as compared with other directions as shown in Figure S5 and S6 (supporting information). Therefore, here the electronic and thermal transport properties for all the samples were evaluated along the sample section perpendicular to the SPS pressurizing direction. The electrical resistivity and Seebeck coefficient were measured in the temperature range ≈300–873 K for Sn1-vS, and in the 300–823 K temperature range for pure SnSe, and undoped and Ag–doped (SnSe)x–(Sn1-vS)1-x samples, respectively, in the helium atmosphere using an electric resistance / Seebeck coefficient measuring system (ZEM–3, Ulvac-Riko, Japan). The total thermal conductivity κT was calculated by the relationship of κT = DCp d where D is the diffusivity, Cp is the specific heat capacity and d is the density of the bulk specimens. d was measured by the Archimedes method as shown in Table S8 (supporting information). D was measured with a laser flash method (NETZSCH Laser Flash Apparatus LFA 457, Germany) as shown in Figure S5 and S9 (supporting information). The specific heat capacity (CP) values of the samples was measured using differential scanning calorimetry thermal analyzer (Netzsch DSC 404 C) technique as displays in Figure S10 (supporting information). Using the Wiedeman–Franz law (κe = LσT) the electronic thermal conductivity (κe) was calculated, where L is the Lorentz, σ is the electrical conductivity and T is the absolute temperature. The lattice thermal conductivities (κL) were calculated according to κL = κT − κe = κT – LσT. Hall coefficient (RH) was measured under a reversible magnetic field of 0.52 T by the Van der Pauw method using a Hall 8

measurement system (ResiTest 8340 DC, Tokyo, Japan) at 300 K. Carrier concentration was calculated using nH = 1/(eRH), and carrier mobility ( µH ) was obtained by µH = σRH .The uncertainty for electrical resistivity was 3%, for the Seebeck coefficient 5% and for thermal conductivity 7%. So, the uncertainty for ZT is 12%.

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ACKNOWLEDGMENT This work was supported by the Basic Science Center Project of NSFC under grant No.51788104 and the National Key R&D Program of China No.2018YFB0703603. Authors Contributions: Jing–Feng Li and Asfandiyar designed the project, Asfandiyar performed part of the experiments and drafted the manuscript, Bowen Cai, Hua–Lu Zhuang and Huaichao Tang helped in characterization and added their input in results & discussions.

Competing Financial Interests statement: The authors declare no competing financial interest.

Tables and Figures

Table 1 Room–temperature (≈300 K) of Hall carrier density and mobility for Sn1-vS, SnSe compounds and SnSe– Sn1-vS solid solutions. All samples are measured along the sample section perpendicular to the SPS pressurizing direction. Samples

nH (1018 cm−3)

µH (cm2/Vs)

SnS

0.00134

4.12

Sn0.99S

0.293

2.47

Sn0.985S

0.344

2.38

Sn0.98S

0.418

1.24

SnSe

0.584

0.642

Sn0.985S0.5Se0.5

0.546

0.457

Sn0.985S0.25Se0.75

1.042

0.386

Sn0.981Ag0.004S0.25Se0.75 Sn0.978Ag0.007S0.25Se0.75 Sn0.975Ag0.01S0.25Se0.75

0.714 0.588 0.438

4.24 2.65 1.83

Table 2.

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Room–temperature (≈300 K) of Hall carrier density and mobility for undoped and Ag–doped Sn0.985S0.25Se0.75 samples. Samples Sn0.985S0.25Se0.75 Sn0.981Ag0.004S0.25Se0.75 Sn0.978Ag0.007S0.25Se0.75 Sn0.975Ag0.01S0.25Se0.75

nH (1018 cm−3) 1.042 0.714 0.588 0.438

µH (cm2/Vs) 0.384 4.24 2.65 1.83

Figure 1. a) Bulk XRD patterns for SnS, Sn0.99S, Sn0.985S, Sn0.98S, Sn0.985S0.5Se0.5 and Sn0.985S0.25Se0.75 samples sintered at 850 and 813 K temperature, respectively and b) Expanded XRD patterns of the Sn0.985S0.25Se0.75, Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 bulk samples showing the presence of AgSnSe2 as a second phase.

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Figure 2. Temperature dependent (a and b) electrical conductivities, (c and d) Seebeck coefficients and (e and f) power factors of the SnS, Sn0.99S, Sn0.985S, Sn0.98S, SnSe, Sn0.985S0.5Se0.5, Sn0.985S0.25Se0.75, Sn0.981Ag0.004S0.25Se0.75, Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 bulk samples along the direction perpendicular to SPS pressurizing direction, respectively.

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Figure 3. Temperature dependent a) total thermal conductivity of the SnS, Sn0.99S, Sn0.985S, Sn0.98S, SnSe, Sn0.985S0.5Se0.5 and Sn0.985S0.25Se0.75 and b) lattice thermal conductivity of the Sn0.985S0.25Se0.75, Sn0.981Ag0.004S0.25Se0.75, Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 bulk samples (inset is the total thermal conductivity of the doped samples).

Figure 4. Temperature dependent ZTs (figure of merit) a) of the SnS, Sn0.985S, SnSe, Sn0.985S0.25Se0.75, Sn0.981Ag0.004S0.25Se0.75, Sn0.978Ag0.007S0.25Se0.75 and Sn0.975Ag0.01S0.25Se0.75 bulk samples and b) Comparisons of the ZT values in this work with other reported values.

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Figure 5. a) HAADF–STEM image of the Sn0.985S sample, b) zoom-in image of the framed area in Figure 5a (solid white), confirm nanoscale dislocations, c) selected area diffraction pattern (SAED) of the matrix, and d) HAADF–STEM images taken inside the grain of Sn0.985S0.25Se0.75 sample, which proves lattice distortions.

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Figure 6. a) STEM annular bright field (ABF) image showing nanoscale precipitates with bright contrast in Sn0.978Ag0.007S0.25Se0.75 sample. The length of the scale bar 500 nm was selected, b) is the EDS elemental map images of the framed area in matrix (dashed orange), c) HAADF–STEM image of the matrix containing one spherical nanoprecipitate, reveals continuous lattice between matrix and precipitate, d) and e) EDS analyses of the matrix and precipitate of the framed areas (dashed white and solid yellow) in (c), shows that the precipitates are rich in Ag and f) HAADF–STEM image, reveals Ag segregates to point defects, e.g., dislocations at the grain boundaries (righ top inset show Ag nanoparticle at the grain boundary).

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• • •

Sn vacancy leads to a ZT of ≈1.1 at 823 K in Sn0.985S0.25Se0.75 sample via improved power factor and reduced thermal conductivity. Upon Ag doping at the sites of Sn in the corresponding sample, ZT was further enhanced to a maximum value ≈1.75. Nanoscale precipitates and dislocations were observed in the high-ZT Ag-doped samples.

Declaration of Interest Statement: The authors declare no competing financial interest.

Biography

First author

Asfandiyar is a Ph.D. candidate in the School of Materials Science and Engineering, Tsinghua University, China. He obtained his MSc. in Physics from Islamia College University of Peshawar (Pakistan) in 2010 and BSc. in Physics from Peshawar University (Pakistan) in 2008. His current research interests include SnS-based thermoelectric materials.

Coauthors

Bowen Cai is now a Postdoctoral fellow working at the School of Materials Science and Engineering in Tsinghua University. His current research interests include thermoelectric materials and high pressure technology. He holds a B.E. from Xi’an University of Technology, a Ph.D from Yanshan University in Qinhuangdao, China.

Hua-Lu Zhuang is a Ph.D. candidate in the School of Materials Science and Engineering, Tsinghua University. He received his Bachelor from Huazhong University of Science and Technology (China) in 2017. His current research focuses on the synthesis of bismuth telluride and their applications in thermoelectrics.

Huaichao Tang is a PhD candidate in the School of Materials Science and Engineering, Tsinghua University. His current research focuses on the synthesis of sulfide-based composites and their applications in thermoelectric.

JingJing-Feng Li is a professor in Tsinghua University, China. He graduated from Huazhong University of Science and Technology (China) in 1984, and obtained his doctor degree from Tohoku University (Japan) in 1991. After working in Tohoku University as assistant professor and associate professor from1992 to 2002, he joined Tsinghua University as a full professor in 2002. His research interests include thermoelectric materials and devices, piezoelectric ceramics, composites and films for MEMS applications.