Ge‐Doped ZnSb/β‐Zn4Sb3 Nanocomposites with High Thermoelectric Performance

ZnSb/β‐Zn4Sb3 nanocomposites are produced from Zn1.1−xGexSb mixtures using a two‐step process. First, proper amounts of the three elements are mixed, melted, and reacted at 800 K. During this process, the nonstoichiometric mixture is crystallized in a combination of ZnSb and β‐Zn4Sb3 phases. Then, the material is ball milled and subsequently hot pressed. Through this process, a dense ZnSb/β‐Zn4Sb3 composite, consisting of β‐Zn4Sb3 nanoinclusions embedded within a ZnSb matrix, is formed. The particular phase distribution of the final ZnSb/β‐Zn4Sb3 composites is a consequence of the harder and more brittle nature of ZnSb than Zn4Sb3, which translates into a stronger reduction of the size of the ZnSb crystal domains during ball milling. This small particle size and the high temperature generated during ball milling result in the melting of the ZnSb phase and the posterior crystallization of the two phases in a ZnSb/β‐Zn4Sb3 matrix/nanoinclusion‐type phase distribution. This particular phase distribution and the presence of Ge result in excellent thermoelectric performances, with power factors up to 1.5 mW m−1 K−2, lattice thermal conductivities down to 0.74 W m−1 K−1, and a thermoelectric figures of merit, ZT, up to 1.2 at 650 K, which is among the highest ZT values reported for ZnSb.


Introduction
Intermetallic compounds with wide compositional ranges and a broad variety of crystal structures provide exciting platforms for both basic research and technological application. In particular, the Zn-Sb binary system displays several intermetallic phases that show complex defect structures and are characterized by excellent transport properties. Several phases of this system have been proposed for thermoelectric (TE) application as alternatives to materials based on expensive and rare elements such as Bi-Te alloys. These Zn-Sb phases include Zn 1+δ Sb, [1,2] β-Zn 8 Sb 7 , [3] Zn 9−δ Sb 7 , [4] α-Zn 3 Sb 2 , [5] Zn 3−x Sb 2 , [6] ZnSb, and β-Zn 4 Sb 3 . Zn 1+δ Sb, which has been synthesized just in nanocrystal form, was predicted to crystallize in a P1 space group consisting of 10 Sb 4− dimers, 8 isolated Sb 3− anions, and 32 Zn 2+ cations. [1] β-Zn 8 Sb 7 , a new metastable phase with Zn 8 Sb 7 composition, has been synthesized via high temperature solid state reaction followed by quenching. β-Zn 8 Sb 7 crystallizes in a non-centrosymmetric orthorhombic space group Pmn21 (No. 31) with unit cell parameters a = 15.029(1) Å, b = 7.7310(5) Å, and c = 12.7431(9) Å. [3] Within the Zn 9−δ Sb 7 subsystem, an α polymorph and a Znpoorer β polymorph with the same hexagonal structure (P6/mmm) but different lattice parameters exist in the temperature range from 514 to 581 °C. The atomic arrangement in Zn 9−δ Sb 7 consists of 3 Sb 2− dimers and 4 isolated Sb 3− ions with 9 Zn 2+ cations. [4] α-Zn 3 Sb 2 quenches in a big hexagonal cell in space group R3 (a = 15.212 (2), c = 74.83(2) Å), where isolated Sb 3− anions construct tetrahedral voids filled with highly disordered Zn 2+ cations. [5] The chemical composition of Zn 3−x Sb 2 is (Zn 2+ ) 17+α (Sb 3− ) 10+2α (Sb 2− ) 2−2α , where α (<1) is the fraction of unpaired Sb 6 atoms. Generally, the composition can be written as Zn 17+α Sb 12 , which is distinct from the Zn 3 Sb 2 phase that contains only Sb 3− anions. [6] All these phases are stable in just a relatively narrow range of temperatures well above ambient, and in the best case scenario they are metastable at room temperature. This instability inhibits their practical application. [3][4][5] β-Zn 4 Sb 3 exhibits the highest TE performance within the Zn-Sb system, with a TE figure of merit, ZT, up to 1.4 at 748 K. [7] However, its moderate stability is a major unsolved issue. β-Zn 4 Sb 3 can decompose to Zn and ZnSb at temperatures below 500 K, what hinders its practical application. [7][8][9] The high ZnSb/β-Zn 4 Sb 3 nanocomposites are produced from Zn 1.1−x Ge x Sb mixtures using a two-step process. First, proper amounts of the three elements are mixed, melted, and reacted at 800 K. During this process, the nonstoichiometric mixture is crystallized in a combination of ZnSb and β-Zn 4 Sb 3 phases. Then, the material is ball milled and subsequently hot pressed. Through this process, a dense ZnSb/β-Zn 4 Sb 3 composite, consisting of β-Zn 4 Sb 3 nanoinclusions embedded within a ZnSb matrix, is formed. The particular phase distribution of the final ZnSb/β-Zn 4 Sb 3 composites is a consequence of the harder and more brittle nature of ZnSb than Zn 4 Sb 3 , which translates into a stronger reduction of the size of the ZnSb crystal domains during ball milling. This small particle size and the high temperature generated during ball milling result in the melting of the ZnSb phase and the posterior crystallization of the two phases in a ZnSb/β-Zn 4 Sb 3 matrix/ nanoinclusion-type phase distribution. This particular phase distribution and the presence of Ge result in excellent thermoelectric performances, with power factors up to 1.5 mW m −1 K −2 , lattice thermal conductivities down to 0.74 W m −1 K −1 , and a thermoelectric figures of merit, ZT, up to 1.2 at 650 K, which is among the highest ZT values reported for ZnSb.
ZT values reached for this material are associated with the complex hexagonal crystal structure of β-Zn 4 Sb 3 with space group 3 R c, which results in a very low thermal conductivity. [10] The crystal structure of β-Zn 4 Sb 3 consists of 6 Sb 2 4− dimers and 18 isolated Sb 3− anions, which need 39 Zn 2+ cations to maintain the charge balance. This structure provides β-Zn 4 Sb 3 with a glass-like thermal conductivity mainly due to the anharmonic motion of the Sb1 atoms coordinated only by Zn atoms. [11] It is also agreed that both interstitial and disordered Zn sites further reduce the already low thermal conductivity of β-Zn 4 Sb 3 . [12] Taking advantage of this low thermal conductivity and a significant Seebeck coefficient, several elements have been introduced in small quantities to optimize the charge carrier concentration of this material and to maximize in this way its ZT, [13][14][15] e.g., Al, [16] Ga, [16] In, [16,17] Hg, [18] Cd, [19] Nb, [20] Te, [21,22] Mg, [23] Ag, [12] Cu, [24] Fe, [25] Bi, [26] Se, [27] Pb, [28] and Ge. [29] Numerous studies have been also devoted to study the thermal decomposition of this phase, its major issue, and to improve it through doping with elements such as Pb, Cd, and Ag. [12,19,28] However, these progresses have not yet meet all the requirements for practical applications, and improving stability remains as a major challenge to be solved for this material.
From a practical point of view, ZnSb not only provides the second highest TE performances within the Zn-Sb system, but it is also the only phase in this system with a proper stability for real TE application. ZnSb crystalizes in an orthorhombic crystal structure (space group Pbca), which contains 8 Zn 2+ cations and 4 Sb 24− dimers per unit cell. ZnSb possesses higher power factors (PFs) than β-Zn 4 Sb 3 , but its relatively higher thermal conductivity, related to its simpler crystal structure, leads to overall lower TE performance. [30] A common approach to reduce the thermal conductivity of this material is through grain refinement to nanoscale dimension. [31][32][33] Okamura et al. [34] synthesized ZnSb by the melting approach and used mechanical grinding as a postsynthesis treatment to reduce its thermal conductivity. Using this strategy, they reduced the thermal conductivity to 1.41 W m −1 K −1 at room temperature and reached a maximum ZT value of 0.8 at 550 K. Thermal conductivities below 1 W m −1 K −1 have been also reported for nanostructured ZnSb prepared by cryomilling and hot pressing. [35] An alternative approach to decrease the thermal conductivity of ZnSb is the incorporation of a second phase in the form of nanoinclusions. This strategy was used in the work of Valset et al., [36] where a ZT value close to unity at 550 K was realized in Cudoped ZnSb containing 2.5 at% of Zn 3 P 2 nanoparticles. The high ZT values reached for this nanocomposite were attributed to a 15% reduction in the thermal conductivity and a modulation doping effect. The same strategy was also used in ZnSb doped with Ag (0.15 at%), Cd (3 at%), and Sn (3 at%), resulting in ZT = 1 at 630 K. [37] The highest ZT value for ZnSb was reported by Xiong et al. [14] in Ag-doped Zn 1−x Ag x Sb (x = 0.002), which reached ZT = 1.15 at 570 K. Being ZnSb the most stable intermetallic phase in the Zn-Sb binary system, to further explore this system through alternative doping elements and phases is a worth endeavor. In this scenario, one potentially interesting dopant that is yet to be systematically investigated is Ge.
In this work, we study the in situ preparation of ZnSb/β-Zn 4 Sb 3 nanocomposites through ball milling off-stoichiometric Zn-Sb mixtures. We additionally analyze the effects of Ge doping on the TE performance of the obtained ZnSb/β-Zn 4 Sb 3 composites.

Phase and Structure
Ge-doped ZnSb/β-Zn 3 Sb 4 nanocomposites with nominal compositions Zn 1.1−x Ge x Sb (x = 0, 0.01, 0.02, 0.04, 0.06) were produced in a two-step process (see details in the Experimental Section). First, proper amounts of the three elements were reacted at high temperature. In a second step, ingots were ground in an agate mortar and placed into a zirconia jar to ball mill them for 20 h using tungsten carbide balls. Ball milled powders were loaded into a 8 mm diameter graphite die and hot pressed under argon atmosphere at 400 °C for 5 min, under a pressure of 100 MPa. The obtained pellets were finally annealed at 380 °C for 1 h under Ar atmosphere.
Using this procedure, Zn 1.1−x Ge x Sb compounds with relative densities ≈95% were obtained. The crystal phase of hot-pressed Zn 1.1−x Ge x Sb nanomaterials was resolved by X-ray diffraction (XRD) (Figure 1). The unsubstituted Zn 1.1 Sb sample was composed of ZnSb and β-Zn 4 Sb 3 phases (JCPDS cards 96-900-8883 and 96-400-1475). Rietveld refinement analysis using GSAS-II software ( Figure S1, Supporting Information) allowed determining the phase fractions of the Zn 1.1 Sb sample to be 59.7% ZnSb and 40.3% β-Zn 4 Sb 3 . No XRD peaks corresponding to unreacted Zn or Sb were observed.
For all Ge-substituted samples, XRD data were also consistent with a mixture of ZnSb and β-Zn 4 Sb 3 phases. As the Ge content increased, XRD patterns more clearly resembled that of pure ZnSb phase. At high Ge substitutional levels, the excess of Zn to Adv. Mater. Interfaces 2019, 6,1900467

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form β-Zn 4 Sb 3 was strongly decreased and just a minor amount of β-Zn 4 Sb 3 could be detected in the XRD patterns. Quantitative Rietveld analyses indicated that the amount of β phase decreased from 40.3% for the unsubstituted sample (x = 0) to 14% for x = 0.06. Besides, a weak peak corresponding to elemental Ge could be already detected in XRD patterns of samples with x = 0.01, and increasingly higher amounts of elemental Ge were observed for substitution levels of x = 0.02 and above.

Microstructural Characterization
After 20 h of ball milling, nanopowders consisted of agglomerated nanoparticles with sizes between 10 and 30 nm. A high tendency to agglomeration was systematically observed and despite long-time sonication during preparation of transmission electron microscopy (TEM) samples, individual particles were rarely observed during TEM studies. Hot-pressed and annealed pellets after manual grinding for TEM and scanning electron microscope (SEM) sample preparation presented irregular shaped particles with different sizes ( Figure S2, Supporting Information). Both, coarse grains and nanodomains with sizes between 50 and 500 nm were observed, indicating clear growth of the agglomerated nanoparticles during hot pressing and posterior annealing. Identical morphological characteristics were observed for all samples, regardless of doping content ( Figure S2b,c, Supporting Information).
High-resolution TEM (HRTEM) characterization of Zn 1.09 Ge 0.01 Sb composites showed them to contain different crystallographic domains (Figure 2). Power spectrum analysis of the different regions observed in HRTEM micrographs allowed discerning two different crystal phases. In Figure 2, the power spectrum of region 1 (R1) revealed the selected crystal structure to match with the Zn 3.878 Sb 3 hexagonal phase (space group = R3-CH) with a = b = 12.2406 Å and c = 12.4361 Å). From this particular HRTEM micrograph, the Zn 3.878 Sb 3 lattice fringe distances were measured to be 0.249, 0.202, and 0.299 nm, at 41.99° and 96.10°, which we interpreted as the Zn 3.878 Sb 3 phase visualized along its [11][12][13][14][15][16][17][18][19][20] zone axis. For the rectangular area in region 2 (R2), the lattice fringe distances were measured to be 0.194, 0.209, and 0.132 nm, at 50.92° and 83.59°, which we interpreted as the orthorhombic ZnSb phase (space group = Pbca) with a = 6.2016 Å, b = 7.7416 Å, and c = 8.0995 Å, visualized along its [−121] zone axis. The comparison between the experimental and the theoretical bulk plane spacing and angles is detailed in Table S1 (Supporting Information). Overall, extensive HRTEM analysis indicated that β-Zn 4 Sb 3 crystal domains were usually surrounded by the ZnSb phase, with different phase ratios depending on the initial Zn-Sb stoichiometry. HRTEM characterization also demonstrated the poly/nanocrystalline nature of the composite and the high crystallinity of the different domains.
The compositional distribution within the nanocomposite was further evaluated by high angle annular dark-field (HAADF)-scanning TEM (STEM) and electron energy loss spectroscopy (EELS) chemical composition maps (Figure 3). EELS compositional maps confirmed the presence of the three elements, Zn, Sb, and Ge, all through the nanopowder. While Ge was distributed homogeneously through the sample, a slightly uneven distribution of Zn and Sb was observed. Most crystal domains of the Zn 1.09 Ge 0.01 Sb sample displayed Sb-rich shells and Zn-rich cores. Taking into account the XRD and HRTEM results displaying the existence of both the ZnSb and β-Zn 4 Sb 3 phases within the sample and the absence of elemental Zn and Sb phases, we attribute the Sb-rich shell to the ZnSb phase and the Zn-rich internal core to the β-Zn 4 Sb 3 phase.
Considering the above results, we visualize the phase distribution within the nanocomposite as consisting in β-Zn 4 Sb 3 grains within a ZnSb matrix. We hypothesize the formation process of such ZnSb/β-Zn 4 Sb 3 nanocomposite as schematized in Figure 3b. During the high energy ball-milling process, the presynthesized Zn 1−x Ge x Sb bulky sample composed of β-Zn 4 Sb 3 and ZnSb phases ( Figure S3, Supporting Information) is shattered by the severe impact of tungsten carbide balls. During ball milling, the ZnSb refines to nanoscale more rapidly than β-Zn 4 Sb 3 because of its higher hardness and lower impact

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toughness. [38][39][40] ZnSb nanocrystals may partially melt during ball milling and its fluid nature may facilitate the coating of the larger β-Zn 4 Sb 3 particles to form β-Zn 4 Sb 3 @ZnSb core@shell structures. Such core-shell structures were found homogeneously dispersed within the remaining ZnSb matrix forming the ZnSb/β-Zn 4 Sb 3 nanocomposite. XRD analysis of presynthesized samples confirmed that ball milling acted as a medium for redistribution of the two phases without substantial structural changes ( Figure S3, Supporting Information).

Thermoelectric Properties
The temperature dependence of the electrical resistivity, Seebeck coefficient, and power factor of Zn 1.1−x Ge x Sb nanocomposites containing different Ge substitution levels is plotted in Figure 4. The electrical resistivity of Zn 1.1 Sb decreased with increasing temperature, from 71.80 to 52.97 µΩm in the temperature range from ambient to 655 K, indicating a nondegenerate semiconductor behavior. The measured room temperature resistivity was close to the value of ≈65 µΩm reported by Xiong et al., [14] but it was considerably lower than values of 200 and 250 µΩm reported by Pothin et al. [30] and Valset et al. [36] for pure ZnSb. The relatively wide range of resistivities reported may be ascribed to the distinct level of intrinsic defects and offstoichiometries and to the different amount and distribution of the β-Zn 4 Sb 3 and ZnSb phases within materials produced using different synthesis techniques. Resistivities measured in the present work were higher than those previously reported for β-Zn 4 Sb 3 compounds, which exhibited degenerate transport properties. Our results suggested that the electrical resistivity of Zn 1.1−x Ge x Sb nanocomposites was a combination of the resistivity of the ZnSb matrix and the β-Zn 4 Sb 3 domains, the presence of which reduced electrical resistivity. Besides, the electrical resistivity significantly decreased when introducing small amounts of Ge. Zn 1.1−x Ge x Sb nanocomposites exhibited p-type conductivity in the entire Ge-substitution range. The Seebeck coefficient decreased with increasing Ge substitution level, following the same trend as the electrical resistivity. The measured decrease of resistivity and Seebeck coefficient with the introduction of Ge may be related to an increase of the charge carrier concentration indirectly associated with the different electronegativity between Ge (2.01) and Zn (1.65) or their slightly different atomic radiuses, Ge (125 pm) and Zn (135 pm), which may result in a larger density of acceptor defects or a lower density of electron donors such as interstitial Zn due to a small lattice contraction. [29] Power factors significantly increased for Adv. Mater. Interfaces 2019, 6,1900467  all Ge-substituted nanocomposites with respect to undoped Zn 1.1 Sb as a result of a high reduction in the electrical resistivity and a moderate decrease in the Seebeck coefficients. All Ge-substituted nanocomposites exhibited very similar power factor in the entire temperature range, while Zn 1.08 Ge 0.02 Sb exhibited a slightly higher PF at high temperatures. The maximum value of power factor for unsubstituted Zn 1.1 Sb samples was 0.89 mW m −1 K −2 at 650 K, which increased for Ge-substituted Zn 1.1−x Ge x Sb nanocomposite up to 1.6 mW m −1 K −2 for x = 0.01 and 1.65 mW m −1 K −2 for x = 0.02 at the same temperature. Figure 5 shows the temperature dependence of the electronic thermal conductivity, κ e , lattice thermal conductivity, κ l , and total thermal conductivity, κ, of Zn 1.1−x Ge x Sb nanocomposites. A relatively low thermal conductivity was measured Adv. Mater. Interfaces 2019, 6,1900467

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Adv. Mater. Interfaces 2019, 6, 1900467 for unsubstituted Zn 1.1 Sb nanocomposites over the entire temperature range measured, reaching down to 0.74 W m −1 K −1 at 475 K. This value is much lower than those usually obtained for ZnSb. These results indicated that the presence of β-Zn 4 Sb 3 / ZnSb interfaces in Zn 1.1−x Ge x Sb nanocomposites strongly reduced the thermal conductivity of this material. The thermal conductivities of Zn 1.1−x Ge x Sb nanocomposites increased with the introduction of Ge. This experimental evidence was explained in part by the increase of the electronic component of the thermal conductivity, and in part by the decrease of the relative amount of β-Zn 4 Sb 3 , which is characterized by an ultralow thermal conductivity. [11] The electronic thermal conductivity was evaluated using the Wiedemann-Franz equation, κ e = LσT, where T is absolute temperature, σ is electrical conductivity, and L is Lorenz number that was estimated using the formula proposed by Snyder and co-workers. [41] The lattice thermal conductivity was then calculated by subtracting the electronic thermal conductivity (κ e ) from the total thermal conductivity (κ t ). As can be seen in Figure 5a, κ e significantly increased with Ge substitution level, especially in the high temperature range measured. For unsubstituted nanocomposite, the value of κ e was 0.079 W m −1 K −1 at room temperature and increased up to 0.20 W m −1 K −1 . On the other hand, κ e was 0.14 W m −1 K −1 at room temperature and increased to 0.44 W m −1 K −1 at 650 K for Zn 1.04 Ge 0.06 Sb. An almost opposite behavior was observed for κ L (Figure 5b), which showed a significant increase with increasing Ge-substitution level at low temperature ranges and an almost Ge-independent behavior at high temperatures. A lattice thermal conductivity of 0.86 W m −1 K −1 was obtained for unsubstituted Zn 1.1 Sb nanocomposites at room temperature, which increased upon Ge doping and reached up to 1.59 W m −1 K −1 for Zn 1.04 Ge 0.06 Sb. Figure 6 shows the temperature dependence of ZT for Zn 1.1−x Ge x Sb nanocomposites. Ge substitution induced an ample increase in ZT as a result of the substantial improvement in the power factor. An exceptional ZT value up to 1.2 was realized in Zn 1.09 Ge 0.01 Sb nanocomposites at 650 K. This high ZT is almost twice the value obtained for unsubstituted Zn 1.1 Sb nanocomposites. Figure 7 shows the temperature dependence of ZT for our best performing Zn 1.09 Ge 0.01 Sb nanocomposite in comparison to those reported in literatures for ZnSb compounds. It can be seen that Zn 1.09 Ge 0.01 Sb has a higher ZT value than all those previously reported values for ZnSb. Especially, the maximum ZT value of 1.2, realized at a substitution level of x = 0.01, was more than twice the value reported for pure ZnSb. [30] To gain a deeper understanding of the TE performance of our nanocomposite, a comparison of thermal conductivity, Seebeck coefficient, and electrical resistivity with previous papers is presented in Figure S4 (Supporting Information). The transport properties of Zn 1.08 Ge 0.01 Sb nanocomposites were comparable to those of Ag- [12] and Cu-doped [36] samples. Our materials exhibited significantly lower thermal conductivities, very close to those reported for ZnSb/SiC nanocomposite, [41] suggesting the dominance of interface scattering. Furthermore, the obtained ZT values for Zn 1.09 Ge 0.01 Sb were comparable to the best performing β-Zn 4 Sb 3 TE materials.
Besides, our ZnSb/β-Zn 4 Sb 3 nanocomposites exhibited superior stabilities than β-Zn 4 Sb 3 TE materials. Zn 1.08 Ge 0.01 Sb nanocomposites exhibited almost identical transport properties even after five cycles of measurements at a temperature up to 650 K ( Figure S5, Supporting Information). XRD patterns of Zn 1.08 Ge 0.01 Sb nanocomposite ( Figure S6, Supporting Information) revealed very minor changes in the relative intensities of ZnSb and β-Zn 4 Sb 3 peaks after five cycles of TE measurements. Quantitative Rietveld analyses indicated that the amount of β phase decreased from 35% to 31% with the TE measurements. Considering the high stability of the ZnSb phase, we believe

Conclusion
We developed ZnSb/β-Zn 4 Sb 3 nanocomposites using a twostep approach. In first step, Zn 1.1−x Ge x Sb (x = 0, 0.01, 0.02, 0.04, 0.06) compounds were produced by a melting approach. In second step, these compounds were ball-milled and hot-pressed. From this procedure, ZnSb/β-Zn 4 Sb 3 nanocomposites consisting on small β-Zn 4 Sb 3 grains surrounded by a ZnSb matrix were obtained. Ge doping led to a substantial decrease in the electrical resistivity concomitant to a moderate decrease in the Seebeck coefficient. As a result, the power factor significantly increased above 1.5 mW m −1 K −2 at 550 K for all doped samples. The high density of ZnSb/β-Zn 4 Sb 3 interfaces resulted in low thermal conductivities, down to 0.74 W m −1 K −1 . Finally, a high and stable ZT value of 1.2 at 650 K was obtained for Zn 1.1−x Ge x Sb at x = 0.01, which is among the highest values reported for ZnSb.

Experimental Section
Sample Preparation: Ge-doped ZnSb/β-Zn 3 Sb 4 nanocomposites with nominal compositions Zn 1.1−x Ge x Sb (x = 0, 0.01, 0.02, 0.04, 0.06) were produced in a two-step process. First, proper amounts of the three elements were reacted at high temperature. High purity Zn (99.9%, Alfa-Aesar), Sb (99.5%, Alfa-Aesar), and Ge (99.999%, Sigma-Aldrich) powders in the proper ratio, Zn 1.1−x Ge x Sb (x = 0, 0.01, 0.02, 0.04, 0.06), were weighted inside an argon-filled glove box and loaded into a quartz ampoule. The ampoule was then repeatedly evacuated and refilled with nitrogen for three times and finally sealed under vacuum, at about 10 −4 Torr. The sealed ampoules were fixed vertically inside a furnace and heated to 527 °C at 100 °C h −1 and soaked at this temperature for 18 h. The furnace was subsequently turned off and the samples were allowed to cool down to room temperature inside the furnace.
In second step, ingots were ground in an agate mortar and placed into a zirconia jar to ball mill them for 20 h using tungsten carbide balls. A ball to powder mass ratio of 6:1 and a rotation speed of 400 rpm were used. The ball milled powders were loaded into a 8 mm diameter graphite die and hot pressed under argon atmosphere at 400 °C for 5 min, under a pressure of 100 MPa. After these 5 min, the heater was turned off and the pressure was kept constant until temperature dropped to about 250 °C, at which point the pressure was released. The obtained pellets were then annealed at 380 °C for 1 h under Ar atmosphere.
Characterization: XRD patterns were collected on a Bruker AXS D8 ADVANCE X-ray diffractometer (Cu-Kα radiation, λ = 0.154 06 Å). The GSAS-II package [42] was utilized for Rietveld refinements analysis. The maximum weighted factor (wR) was 4.43% for Zn 1.1 Sb and the minimum wR was 3.90% for Zn 1.1−x Ge x Sb (x = 0.06) at room temperature. TEM studies were performed on a Zeiss Libra 120, operated at 120 kV. The morphology of the samples was examined using a Zeiss Auriga SEM at 5.0 kV. Energy dispersive X-ray spectroscopy, worked with Zeiss Auriga SEM device, was employed for elemental analysis. HRTEM and STEM studies were carried out using a field emission gun FEI Tecnai F20 microscope at 200 kV with a point-to-point resolution of 0.19 nm. HAADF-STEM was combined with EELS in the Tecnai microscope by using a GATAN QUANTUM filter. Seebeck coefficient and electrical resistivity were measured simultaneously during the sample heating up in a LSR-3 LINSEIS system under helium atmosphere. Thermal conductivity was obtained as a product of thermal diffusivity (λ), heat capacity (C p ), and mass density of the specimen (ρ), (k = λ × ρ × C p ). Thermal diffusivity was measured during the sample heating up using a XFA 600 Xenon Flash apparatus. The density values were obtained using the Archimedes' method. The Dulong-Petit approximation was used to obtain the specific heat (C p ).

Supporting Information
Supporting Information is available from the Wiley Online Library or from the author.