Thermoelectric materials are becoming more and more concerned, aiming at boosting their applications in waste-heat recovery and cooling.^[1,2] Generally, the thermoelectric conversion efficiency can be determined by a figure of merit , where T, S, σ, and κ are absolute temperature, the Seebeck coefficient, electrical conductivity, and thermal conductivity of a thermoelectric material, respectively. Obviously, high zT is equivalent to high power factor ( and low thermal conductivity simultaneously. In the past decade, several strategies were usually adopted to enhance the power factor of bulk materials, such as carrier concentration engineering,^[3,4] increase of band degeneracy,^[4–6] enhancement of band effective mass,^[7,8] and improvement of carrier mobility.^[9] To decrease the thermal conductivity, which contributes most significant advances in improving zT, there are several principles, including introducing point defects,^[10] second-phase nanostructures,^[11–13] all-length-scale hierarchical architectures,^[14,15] and choosing compounds with complex crystal structures.^[16]

In 2014, Zhao et al. reported that a record zT ∼2.6 at 923 K was achieved along the b axis of SnSe single crystal, benefiting from the low lattice thermal conductivity ( at 973 K) which was mainly ascribed to large phonon anharmonicity.^[17,18] Furthermore, enhanced power factor and ultrahigh average zT (1.60) in the range from 300 K to 923 K were realized in Na-doped samples.^[4,6,19] These exciting phenomena have stimulated much interest on this selenide which consists of earth-abundant and environmentally friendly elements. However, the application of thermoelectric SnSe single crystals is hindered because their growth is challenging due to a strong phase transition at high temperature and machinability is poor owing to the weak chemical bonds along a axis. As a result, considerable effort has been focused on improving the thermoelectric performance of polycrystalline SnSe.^[20–25] Unfortunately, up to date, the ultrahigh zT values of SnSe crystals have not been reproduced in polycrystalline SnSe because polycrystalline samples show lower carrier mobility and higher thermal conductivity. It is reasonable to suppose that the carrier scattering at grain boundaries and random distribution of crystalline grains with strong anisotropy should decrease the carrier mobility. Unexpectedly, the higher lattice thermal conductivity of polycrystalline SnSe is reported by lots of papers, which is at odds with the simple expectation that grain boundary scattering should act as an additional source of phonon scattering.^[26] The possible oxidation in polycrystalline SnSe could reconcile this contradicting result since SnO has very high thermal conductivity ( ).^[27,28] Besides, vast off-stoichiometric defects^[10] and relatively low density^[29] may provide additional contribution to the ultralow lattice thermal conductivity.

To improve the thermoelectric performance of polycrystalline SnSe, Na,^[30] K,^[21,31] Ag,^[22] Br,^[23] I,^[24] and BiCl₃^[25] have all been proven to be effective dopants in both p-type and n-type samples. Nevertheless, the maximum zT nearly keeps at relatively low values ( ) though the carrier concentration is effectively increased. SnSe crystallizes in a highly anisotropic layered orthorhombic (Pnma group) crystal structure below 800 K and undergoes a structural transition from Pnma to Cmcm above 800 K.^[32,33] The layered structure suggests that the control of crystalline orientation is very significant in polycrystalline SnSe samples.^[34,35] Besides, Sn vacancy always plays an important role in improving the thermoelectric properties because it resembles hole-doping which can enhance the carrier concentration of p-type SnSe by shifting the Fermi level and simultaneously inducing localized flat bands in the vicinity of the Fermi energy level.^[36–38] Meanwhile, it also acts as point defect to scatter phonon and then decrease the thermal conductivity. These motivate us to optimize the thermoelectric properties by simultaneously regulating the anisotropic crystal growth and Sn vacancy.

In this study, the SnSe polycrystalline bulk materials were successfully prepared by heat treating the SnSe₂ nanostructures synthesized via a wet-chemistry route at temperature above 600 °C. SnSe₂ is a semiconductor with layered CdI₂-type structure, and has very strong anisotropic electrical and thermal properties.^[39–42] This distinctive crystalline orientation of precursor SnSe₂ may affect the anisotropic crystal growth during the transformation of SnSe₂ to SnSe. Furthermore, vast Sn vacancy should be formed during this process because of the sublimation of atoms. In fact, the measurement of thermoelectric properties demonstrate that the intensified anisotropic crystal growth and Sn vacancy enhance the power factor from to at 612 °C. Then, a high zT value of 1.07 at 612 °C is obtained.

In this experiments, the temperature of the condensing reflux method in solvent was regulated. After the reaction, black powders were obtained. The corresponding XRD patterns are shown in Fig. 1(a). From this figure, we could conclude that only peaks well-matched with Se (JCPDS 01-086-2246) are observed if the temperature does not exceed 100 °C. However, the XRD patterns of samples synthesized at temperatures above 100 °C could be indexed to SnSe₂ phase (JCPDS 01-089-2939) without any impure phase. Furthermore, the intensity of peaks becomes stronger and the corresponding full-width at half-maximum (FWHM) decreases with increasing reaction temperature. This result suggests that the elevated temperature activates the reaction and then increases the grain size of SnSe₂ nanostructures. Similarly, it is unable to obtain SnSe nanostructures at temperature below 110 °C as reflected by Fig. 1(b). Fortunately, the XRD patterns of powders obtained at the temperature above 110 °C could be ascribed to SnSe crystal structure (JCPDS 00-048-1224). No obvious peaks of impurities can be observed when the reaction temperature exceeds 140 °C, demonstrating the successful synthesis of SnSe nanostructures via this wet-chemistry method.

Fig. 1.

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	Fig. 1. (color online) XRD patterns of (a) SnSe2 and (b) SnSe powders synthesized via a wet-chemistry route at different reaction temperatures.

The morphologies of the SnSe₂ and SnSe nanostructures were observed by SEM and TEM. As given in Fig. 2, the characterization of SEM shows that the SnSe₂ nanostructures obtained at 140 °C have the shape of nano-petal (Fig. 2(a)). On the other hand, the morphology of the sample synthesized at 182 °C looks more like potato chips with nanosize (Fig. 2(b)). Obviously, the increasing of reaction temperature increases the size of nanostructures, which is in accordance with the XRD results. In detail, the width of these two-dimensional nanostructures is expanded from nanometer (140 °C) to micrometer (182 °C). Though SnSe also has layered crystal structure, the SnSe nanostructures display agglomerated nanoparticles. Meanwhile, the sample obtained at 182 °C (Fig. 2(d)) presents larger agglomerations than that of the one prepared at 140 °C (Fig. 2(c)). Figure 3 shows the typical TEM and HRTEM images. In agreement with the observation of SEM, the SnSe₂ sample synthesized at 182 °C (Fig. 3(a)) reveals nanosheets with the thickness of several nanometers to about 50 nm and the width of several hundred nanometers to about . In contrast, nanoparticles with the size of several nanometers to about 25 nm are absorbed on the surface of lumps in the SnSe sample obtained at 182 °C (Fig. 3(b)). The HRTEM characterization reconfirms the crystal structures of SnSe₂ and SnSe by matching the crystal plane spacing as shown in the insets of Figs. 3(a) and 3(b), respectively.

Fig. 2.

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	Fig. 2. SEM images of SnSe2 and SnSe nanostructures synthesized at 140 °C [(a) and (c)] and 182 °C [(b) and (d)], respectively.

Fig. 3.

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	Fig. 3. (color online) TEM images of (a) SnSe2 and (b) SnSe nanostructures synthesized at 182 °C. The insets are their corresponding HRTEM images.

We prepared the polycrystalline SnSe via two methods as described in Section 2. The only difference between them is the choice of precursors, namely SnSe₂ and SnSe nanostructures. The structure of the samples was investigated by XRD measurement. The collected patterns are given in Fig. 4 and demonstrate that the prepared SnSe₂ keeps its original crystal structure until the temperature of heat-treating is elevated to 550 °C, where some peaks indexed to SnSe phase (JCPDS 00-014-0059) appear. With further increasing the temperature, the precursor of SnSe₂ transforms into SnSe phase completely at 600 °C, without observable diffraction of impurities. Due to the higher vapor pressure of Se than that of Sn, it is anticipated that the sublimation of Se should be stronger than that of Sn, driving the formation of SnSe phase at high temperature ( ). When the temperature reaches 850 °C, a small amount of powders were collected in our experiments, which is not enough to fabricate bulk materials. Interestingly, the obtained SnSe phase indicates a strong anisotropic crystal growth, which is reflected by the fact that the diffraction intensity of (400) crystal face is particularly stronger than the others. The bc orientation degree for the (h00) crystal planes termed as F (h00) can be calculated with Lotgering method by the following equations^[42] where P and P₀ are the ratios of the integrated intensities of all (h00) crystal planes to those of all (hkl) planes for preferentially and randomly oriented samples, respectively. F = 0 and F = 1 mean completely disordered and ordered, respectively. In the present work, the P and P₀ are calculated by using the XRD data and the standard card (JCPDS No. 00-014-0159). The result indicates that the values are 0.51, 0.76, and 0.89 for the XRD patterns of the samples with heat-treating temperature of 600 °C, 700 °C, and 800 °C, respectively. This improved anisotropic crystal growth along bc plane should be beneficial to improving the thermoelectric properties since the direction along b axis has the best thermoelectric performance.^[17] Oppositely, the strongest diffraction peak of the samples prepared by using SnSe nanostructures as the precursor could be ascribed to (111) crystal face, which is typically shown in the inset in Fig. 4.

Fig. 4.

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	Fig. 4. (color online) XRD patterns of the polycrystalline obtained by treating the precursors of SnSe2 nanostructures in a tubular furnace at different temperatures. The inset is the XRD patterns of the reference, which prepared by using SnSe nanostructures as the precursor.

The phenomena of anisotropic crystal growth are also affirmed by the SEM observation, which was performed on the fracture surface parallel to the pressing direction of the SPS process. Figure 5(a)–5(c) give the images of the sample 400 °C@SnSe₂, 600 °C@SnSe₂, and 800 °C@SnSe₂, respectively. Apparently, all samples show a layered morphology. Meanwhile, the size of slice becomes larger and larger with increasing the temperature. This is reasonable because that both SnSe₂ and SnSe have a layered structure. However, as shown in Fig. 5(d), the anisotropic crystal growth of the sample 600 °C@SnSe is weaker.

Fig. 5.

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	Fig. 5. (color online) SEM images of the fracture surface parallel to the pressing direction of SPS process: (a) 400 °C@SnSe2, (b) 600 °C@SnSe2, (c) 800 °C@SnSe2, and (d) 600 °C@SnSe.

Figure 6 displays the thermoelectric properties being measured perpendicular to the pressing direction of the SPS process, which present a better thermoelectric performance than those being measured the parallel to the pressing direction in our experiment. In Fig. 6(a), we find that the electrical resistivity increases gradually at temperature ranging from 50 °C to about 300 °C, and then decreases sharply at temperature ranging from 300 °C to about 550 °C for the SnSe bulk materials prepared by using SnSe₂. Meanwhile, the increasing of the heat-treating temperature decreases the electrical resistivity from (600 °C) to (800 °C) at about 612 °C. This phenomenon may be mainly owing to the increased grain size as observed by SEM, which should improve the carrier mobility due to the decreased carrier scattering at grain boundaries. On the other hand, the sample 600 °C@SnSe has higher electrical resistivity. Generally, the electrical conductivity is determined by the carrier concentration and mobility according to the formula . Therefore, anisotropic crystal growth in polycrystalline SnSe along bc plane would favor the transport of carriers because the carrier mobility in SnSe crystals is higher along b and c axes than that of a axis. Based on the observation of SEM, which shows that the sample 600 °C@SnSe has weaker anisotropic crystal growth than the others, it is easy to anticipate that the carrier mobility in sample 600 °C@SnSe would be lower. Besides, it is reasonable to conclude that the sample T@SnSe₂ contains more Sn vacancy than T@SnSe, since the former (1.12/1) has a higher / value than that of the latter one (1.03/1) as shown in Table 1. Usually, Sn vacancy could provide holes,^[34] and then both the increased carrier mobility and concentration in sample T@SnSe₂ would decrease the electrical resistivity. The evaluated Hall concentration and mobility at room temperature as shown in Table 2 also agree well with the above conclusion as follows. The sample 600 °C@SnSe₂ ( , ) owns higher carrier concentration and mobility than the sample 600 °C@SnSe ( , ), suggesting that the precursor of SnSe₂ nanostructures is more beneficial for the generation of Sn vacancy and anisotropic crystal growth. At the same time, the elevated heat-treating temperature may also enhance the formation of Sn vacancies and anisotropic crystal growth since the carrier concentration and mobility are improved to and for the sample 800 °C@SnSe₂, respectively. The Seebeck coefficient measurement given in Fig. 6(b) shows that SnSe₂ bulk materials have negative values, suggesting that they are n-type semiconductors, while polycrystalline SnSe bulk materials presents positive values, reflecting their p-type properties. This is consistent with the results in earlier reports.^[34,40] Besides, the Seebeck coefficient depending on the heat-treating temperature varies limitedly. For example, the Seebeck coefficient of the sample 800 °C@SnSe₂ is while that of the sample 700 °C@SnSe₂ is at 612 °C. In line with the trend of electrical resistivity, the Seebeck coefficient of the sample 600 °C@SnSe is reasonably higher. However, because of its higher electrical resistivity, the sample 600 °C@SnSe shows very low power factor as revealed in Fig. 6(c), only at 537 °C. In the samples T@SnSe₂, thanks to the remarkably improved electrical transport, the power factor is enhanced from to at 612 °C, which is comparable with the single crystal.^[17] The thermal diffusivity and thermal conductivity are shown in Figs. 6(d) and 6(e), respectively. From these figures, it is easy to conclude that the thermal transport in SnSe is more difficult than that in SnSe₂. For instance, in comparison with the sample 800 °C@SnSe₂ (1.32 W/mK), the thermal conductivity of the sample 400 °C@SnSe₂ is almost twice at 51 °C ( ). As shown in the inset of Fig. 6(e), the lattice thermal conductivity of the sample 700 °C@SnSe₂ locates only a little bit below that of the sample 800 °C@SnSe₂ above 300 °C. Meanwhile, the thermal conductivity of the sample 800 °C@SnSe₂ is also relatively low, only at 612 °C. As a result of the greatly improved power factor and relatively low thermal conductivity, a maximum zT value of 1.07 at 612 °C is obtained by the sample 800 °C@SnSe₂. To the best of our knowledge, it is one of the best results of SnSe polycrystalline samples without foreign-atomic doping.

Fig. 6.

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Fig. 6. (color online) Thermoelectric properties of the polycrystalline along the orientation perpendicular to the pressing direction of SPS process: (a) electrical resistivity, (b) Seebeck coefficient, (c) power factor, (d) thermal diffusivity, (e) thermal conductivity ( , and (f) zT. The inset in panel (e) is the lattice thermal conductivity ( , which is calculated based on . The carrier thermal conductivity ( is estimated based on the Wiedemann–Franz law ( ). The Lorenz number is derived from Fermi energy with a simple one-band model and only acoustic phonon scattering considered.

Table 1.

Table 1.

Table 1. The chemical composition of the polycrystalline samples obtained by XRF measurement. . Sample Mass fraction Mass fraction Chemical of Se/% of Sn/% fomula 400 °C@SnSe2 53.75616 41.09509 SnSe1.97 800 °C@SnSe2 41.11609 55.34402 SnSe1.12 600 °C@SnSe 38.94319 56.97995 SnSe1.03

Table 1. The chemical composition of the polycrystalline samples obtained by XRF measurement. .

Table 2.

Table 2.

Table 2. The carrier concentration and Hall mobility of the polycrystalline samples. . Sample Carrier concentration/cm−3 Hall mobility/( ) 800 °C@SnSe2 4.81 × 1018 36.4 600 °C@SnSe2 8.45 × 1017 32.4 600 °C@SnSe 7.02 × 1017 21.6

Table 2. The carrier concentration and Hall mobility of the polycrystalline samples. .

In this study, by using SnSe₂ nanostructures as the precursor, polycrystalline SnSe was successfully prepared by a combined heat-treating process in a reducing atmosphere and following spark plasma sintering procedure It shows better thermoelectric properties than that of the reference sample 600 °C@SnSe. Owing to the reduction of electrical resistivity, the power factor is enhanced from to at 612 °C. XRD, SEM, XRF, and Hall effect characterizations suggest that anisotropic crystal growth and Sn vacancy may be responsible for the enhanced electrical transport. In addition, the thermal conductivity has a weak relationship with the heat-treating temperature As a result, a high zT value of 1.07 at 612 °C is achieved, which is one of the best results of SnSe polycrystalline samples without foreign-atomic doping. This study provides a new method to synthesize polycrystalline SnSe and pave a way to improve the thermoelectric properties of polycrystalline bulk materials with similar layered structure.