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Functional Inorganic Materials and Devices

Modification of bulk heterojunction and Cl doping for high thermoelectric performance SnSe/SnSe nano-composites 2

Yuejiao Shu, Xianli Su, Hongyao Xie, Gang Zheng, Wei Liu, Yonggao Yan, Tingting Luo, Xiao Yang, Dongwang Yang, Ctirad Uher, and Xinfeng Tang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b00524 • Publication Date (Web): 19 Apr 2018 Downloaded from http://pubs.acs.org on April 19, 2018

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ACS Applied Materials & Interfaces

Modification of bulk heterojunction and Cl doping for high thermoelectric performance SnSe2/SnSe nano-composites *

Yuejiao Shu,a Xianli Su,a, Hongyao Xie,a Gang Zheng,a Wei Liu,a Yonggao Yan,a Tingting Luo, a Xiao Yang,a Dongwang Yang,a Ctirad Uher,b Xinfeng Tanga, a

State Key Laboratory of Advanced Technology for Materials Synthesis and

Processing, Wuhan University of Technology, Wuhan 430070, China b

*

*

Department of Physics, University of Michigan, Ann Arbor, Michigan 48109, USA Correspondence to: X. Su ([email protected]), X. Tang ([email protected]).

1

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Abstract: SnSe2 is a wide band gap semiconductor (Eg = 1.05 eV) with a typical two-dimensional hexagonal crystal structure of the prototype CdI2 phase, resulting in an intrinsically low thermal conductivity, which is favorable for thermoelectrics. Herein, we reported the remarkable role of Cl doping in SnSe2/SnSe nanocomposites. Doping with Cl in the system not only increased the carrier concentration by an order of magnitude but it also modified the heterojunction from that of the Schottky junction type (p-n junction) in undoped samples to junctions having an Ohmic contact (n-n junction) when the samples were doped with Cl, increasing their carrier mobility in the process. On account of the simultaneously boosted carrier concentration and carrier mobility upon Cl doping, the electrical conductivity and the power factor were greatly increased. Moreover, the enhanced point defect phonon scattering induced by Cl doping, coupled with the interface phonon scattering, resulted in a suppression of the thermal conductivity. As a consequence, the maximum ZT value of 0.56 at 773 K was achieved in the 6% Cl-doped SnSe2/SnSe nanocomposite measured in the direction parallel to the pressing direction. This is an almost 6 times larger value than measured on the undoped composite. In addition, unlike the conventional layered compounds (Bi2Te3, SnSe), the ZT value parallel to the pressing direction is much higher than the one measured perpendicular to the pressing direction. This study provides a new way for optimizing the thermoelectric properties of materials through interface regulation. Keywords:

SnSe2

based

material,

SnSe

nanoprecipitates,

thermoelectric properties, Cl doping. 2


heterojunction,

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Introduction Thermoelectric (TE) materials can achieve a direct conversion between heat and electricity and vice versa via the Seebeck and Peltier effects, with the advantages of emissions-free, no moving parts, high reliability and long life span.1-5 Therefore, thermoelectricity can play an important role in solving the energy and environmental crisis. The performance of TE materials can be evaluated by the dimensionless thermoelectric figure of merit defined as ZT=Tα2σ/κ, where α, σ, κ and T are the Seebeck coefficient, electrical conductivity, thermal conductivity and absolute temperature, respectively.6,

7

In the past few decades, immense efforts have been

devoted to improving thermoelectric properties and developing high performance TE materials through doping and electronic band structure manipulations, and forming solid solutions and introducing nanostructural features in the crystal lattice to enhance phonon scattering and thus lower the thermal conductivity.8-13 Despite much improved thermoelectric performance of the state-of-the-art thermoelectric materials, such as Bi2Te3,14, 15 CoSb3,16, 17 PbTe,18, 19 GeTe,20, 21 and half-Heusler alloys,22, 23 their large-scale commercial applications have not yet materialized because most of them contain expensive, low abundant and often toxic heavy metal elements. Consequently, it is essential to explore new, more environmentally friendly and cost effective high-performance TE materials in order to be competitive and make an impact on large scale energy conversion applications. Metal dichalcogenides, possessing natural super-lattice structure, are an important class of materials with some unique electronic and optical properties.24 3



Currently, they are attracting considerable attention as 2D tunable semiconductors.25 Among them, tin diselenide (SnSe2) possesses a hexagonal crystal structure with the prototype CdI2 phase and non-toxic low-cost elements. The structure forms two-dimensional atomic layers in the sequence of Se−Sn−Se bounded by van der Waals forces between the inter-layers,26-28 as shown in Figure 1. The above structural characteristics and diverse chemical bonding give SnSe2 compounds highly anisotropic transport properties and an intrinsically low lattice thermal conductivity, both aspects are highly favorable for thermoelectricity.29 Un-doped SnSe2 possesses a wide band gap of 1.05 eV and usually exhibits an n-type semiconducting transport behavior. Accordingly, it shows a high Seebeck coefficient (-450 µV K-1) at room temperature.30 However, the intrinsically low carrier concentration (1017 cm-3) and low carrier mobility (30 cm-2V-1s-1 for single crystal) lead to a relative low electrical conductivity in comparison to other conventional TE materials,31 and result in the inferior thermoelectric properties of un-doped SnSe2. Theoretical calculations demonstrated that excellent thermoelectric properties can be attained through n-type doping with ZT values estimated as high as 2.95 at 773 K along the crystallographic b-axis when the carrier concentration is tuned to 1019 cm-3.32 Thus, n-type SnSe2 is considered a promising thermoelectric material. However, these are theoretical predictions and very few reports describe the actual experimental studies.33-35 Saha et al. synthesized Cl-doped SnSe2 nanosheets in which Cl doping increases the carrier concentration and leads to an enhanced ZT value of 0.22 at 610 K.36 Li et al. used Ag as a dopant in the SnSe2 system to improve its carrier mobility from 1.58 cm-2V-1s-1 in 4


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SnSe2 to 3.54 cm-2V-1s-1 in Ag-doped SnSe2 at room temperature. This improved the figure of merit ZT to 0.4 at 773 K.37 In spite of notable improvements upon doping the SnSe2 structure, the carrier concentration still remains far below the value required to achieve the optimum carrier concentration accompanied with the low carrier mobility. Therefore, increasing the carrier concentration and carrier mobility simultaneously is the key to optimizing thermoelectric properties of SnSe2. It is commonly recognized that forming in-situ nano-composites is an important method to improve the performance of TE materials.38, 39 For one thing, the in-situ formed nanostructures can increase interface phonon scattering and reduce the thermal conductivity.40-42 Moreover, the two phase interface that has formed can scatter electrons and/or holes selectively and reduce transmission of electron-hole pairs, inhibiting the bipolar thermal conductivity when the structure enters the regime where intrinsic excitations dominate.43, 44 In addition, the potential barrier that forms in the presence of a two phase interface can improve the Seebeck coefficient, yet it exerts a very weak effect on the electrical conductivity though energy filtering, whereby the low energy carriers are scattered out leaving behind high energy charge carriers to facilitate charge transport, ultimately improving thermoelectric properties of the material.45-47 Depending on the charge carrier type of the materials forming the two phase interface, the interfaces can be divided into homogeneous junctions (p-p or n-n junctions) and heterojunctions (p-n junctions). Different junction structures have different effects on the charge and phonon transport properties since the space charge distribution at the interface is different.48 There are very few reports to quantitatively 5



account for the influence of different forms of junction on the thermoelectric transport properties, especially as far as concerns bulk thermoelectric materials. In this work, we have successfully prepared SnSe2/SnSe bulk composites with in-situ formed SnSe nanoprecipitates by melting-quenching combined with spark plasma sintering (SPS). The correlation between Cl doping, its effect on the junction structure, and the thermoelectric transport properties was investigated. The p-n junction formed at the interface between SnSe and SnSe2 of the undoped SnSe2/SnSe composite gives rise to a built-in voltage at the interfacial area, which scatters electrons of the matrix and results in a low carrier mobility of the undoped composite. In contrast, doping the composite with Cl alters the interface from the p-n junction character of the undoped SnSe2/SnSe composite to the n-n junction with an ohmic contact, modifying the spatial charge distribution in the process. As a consequence, both the carrier concentration and mobility have significantly improved in Cl-doped composites in comparison to SnSe2/SnSe composites with no doping. In addition, in Cl-doped composites the thermal conductivity has been suppressed due to enhanced point defect phonon scattering and interfacial phonon scattering. A sample with the nominal Cl content of 0.06 attained the maximum ZT value of 0.56 at 773 K when measured in the direction parallel to the direction of the applied pressure during the SPS process. This value is about 6 times larger than what was measured on the undoped composite sample. In addition, unlike the conventional layered compounds (Bi2Te3, SnSe),49-51 the ZT value measured parallel to the pressing direction is much higher than the one measured perpendicular to the pressing direction. 6


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Experimental: High purity powders of Sn (5N), Se (5N) and SnCl2 (4N) were weighed according to the stoichiometric ratio of SnSe2-xClx (x = 0, 0.015, 0.025, 0.035, 0.05, 0.06, and 0.07), and the mixed powders were sealed in evacuated quartz tubes (diameter of 20 mm) and heated slowly up to 973 K. The tubes were held at this temperature for 24 h, and then quenched rapidly to room temperature in cold water. Although we weighed the material according to the stoichiometric ratio, it always results in a SnSe2/SnSe composite with SnSe2 as the majority phase and SnSe as the minority secondary phase. Therefore, in the result and discussion parts, we labeled the samples as x%Cl-SnSe2/SnSe composites which is corresponding to the nominal composition of SnSe2-xClx samples. The obtained ingots were ground into a fine powder, sieved through 200 mesh sieves, and subsequently vacuum sintered using the spark plasma sintering (SPS) apparatus operated under a pressure of 30 MPa at 773 K for 8 min to obtain bulk samples with the relative density over 98%. The phase composition of the samples was identified by powder X-ray diffraction (PANalytical–Empyrean; Cu Kα). The electrical conductivity and the Seebeck coefficient were measured simultaneously using commercial equipment (ZEM-3, Ulvac Riko, Inc.) under a low pressure of inert gas (He). The thermal conductivity was calculated according to the relationship κ =λCpρ, where λ, Cp and

ρ are thermal diffusivity, the heat capacity, and the density of bulk samples, respectively. The thermal diffusivity (λ) was measured using the laser flash system 7



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(LFA 457; Netzsch) in an argon atmosphere. We used the Dulong-Petit limit as the specific heat capacity (Cp). The sample density (ρ) was measured by the Archimedes method. The electrical conductivity and the Hall coefficient from 10 to 300 K were measured in a physical property measurement system (PPMS-9T, Quantum Design). The carrier concentration (n) and the carrier mobility (µH) were calculated by the following equations: n = 1/eRH and µ = σ/ne. The morphology, including elemental distribution mapping of samples, was characterized by field emission scanning electron microscopy (FESEM, Hitachi SU-8020, Japan), and high resolution transmission electron microscopy (HRTEM, JEM-2100F, JEOL, Japan). The chemical valence of elements was determined using XPS (VG Multilab 2000). The work function was measured by using UPS (VG Multilab 2000), and the Fermi level is set to 0 through the calibration of the gold sheet. The orientation factor F was calculated according to the following equation: =

(1)

where P and P0 are the ratios of the integral intensities of the (00l) planes and are calculated by equations: ∑ I(00l) / ∑(hkl), I is the intensity of the diffraction peak of the sample in the XRD diagram.

3. Results and discussion 3.1 Composition and microstructure Powder X-ray diffraction patterns of x%Cl-SnSe2/SnSe composites (0≤x≤7) are shown in Figure 2(a). The diffraction peaks are consistent with the standard 8


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pattern of SnSe2 (JCPDS#00-035-0752). However, a very weak additional peak of the SnSe phase is detected in all samples regardless of the doping content of Cl, which is due to Se element deviates from the stoichiometric ratio easily during the process of synthesis (Figure S1, reproduced with permission from reference 52. Copyright 1998 Springer Nature). 52Actually, the mixture phase has been observed in several previous publications, in spite of that they claimed the single phase SnSe2 is attained.25, 53 The expanded view of XRD in the range of 59.6o to 60.8o is shown in Figure 2(b). Apparently, the diffraction peak shifts to higher angles with the increasing content of Cl, indicating that Cl successfully substitutes on the Se site. The contraction of the lattice following Cl doping is mainly due to a smaller the atomic radius of Cl (1.81 Å) compared to that of Se (1.98 Å). Figure 2(c) displays the XRD pattern of bulk undoped composite after spark plasma sintering along both the pressing direction and perpendicular to the pressing direction. Relative intensities of (00l) diffraction peaks in the direction perpendicular to the pressing direction are much stronger than those along the pressing direction, revealing that the SPS sintering process, indeed, has an important influence on the texture of the samples. As is well known, the crystallographic orientation of two dimensional structures has a strong influence on charge and phonon transport properties. Thus, in order to characterize the degree of orientation of the samples, we calculated their orientation factor F using the XRD data collected on samples oriented perpendicular to the pressing direction along the (001) crystal surface. When samples have no preferred orientation, the orientation factor F equals to 0. For fully oriented samples the orientation factor F is 1. The effect of Cl 9



content on the orientation is shown in Figure 2(d). The orientation factor F is essentially constant and fixed at a value of 0.5 as the content of Cl varies. Doping with Cl therefore has essentially no impact on the orientation of samples. However, because the spark plasma sintering process gives rise to the orientation dependence of samples, the influence of orientation must be taken into account when testing and analyzing their thermoelectric properties. Microstructures on the freshly fractured surfaces of all sintered samples were observed by FESEM. Because the morphologies of all samples are quite similar, we have selected the undoped SnSe2/SnSe composite as a representative case. Figure 3(a) and 3(b) show FESEM images of the freshly fractured surfaces of undoped composite in the direction perpendicular to the pressing direction. The sintered samples are fully condensed with a relative density above 98%. A typical lamellar structure with the coarse grain size distribution is observed, indicating high degree of orientation in the sintered samples, the result consistent with the XRD data. To measure sample homogeneity, backscattered electron (BSE) images of the polished surface and the element distribution maps for sintered 6% Cl-doped composite are shown in Figures 3(c)-3(f). There is no contrast seen in Figure 3(c) that would suggest the presence of impurity phases. In addition, element maps in Figures 3(d)-3(f) document that Sn, Se and Cl are distributed homogeneously on the micro-scale. This further confirms that Cl has effectively substituted on the Se site. Although XRD results show that a SnSe secondary phase has been detected in 10


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all the samples, the presence of such a phase is not observed in the BSE image. This may be due to the smaller grain size in the secondary SnSe phase, which is apparently well-distributed and embedded inside SnSe2 grains. To learn more about the microstructure of bulk samples, High-Resolution Transmission Electron Microscopy (HRTEM) was applied to characterize the microstructure. Figure 4(a) is a typical low magnification TEM image of the SnSe2 sample. The electron diffraction analysis carried out in the area delineated by the dashed white square in Figure 4(a) reveals that the matrix is SnSe2. A dark structure within the large white circle in Figure 4(a) yields two sets of diffraction patterns in the selected area electron diffraction, shown in Figure 4(b). The brighter and the less bright diffraction spots belong to SnSe2, respectively to SnSe as clearly follows from indexing the diffraction spot patterns. For instance, the bright diffraction spots yield inter-planar spacings of 0.29 nm and 0.31 nm that correspond to (011) and (002) planes of SnSe2, while the less intense spots are those associated with inter-planar spacings of 0.30 nm and 0.27 nm that fit better with (011) and (211) planes of SnSe. In addition, the EDS energy spectrum analysis of the large white circular region gives the ratio of Se to Sn as 53.9%:32.7% (Figure S2), which is in the range between that of SnSe2 and SnSe, verifying the presence of SnSe in this region. Figure 4(a) also depicts some gray nano-domains in the sample and their high resolution TEM is shown in Figures 4(c)-4(d). A large layered grain is shown in Figure 4(c) with the interplanar spacing of 0.33 nm, corresponding to the (010) plane of SnSe2. 11



Also, there is a small SnSe nano-grain on the left with the interplanar spacing of 0.293 nm and 0.303 nm, corresponding to the (111) and (011) planes of SnSe, respectively. Figure 4(d) displays a SnSe nanoparticle with the size of 10-15 nm embedded in the SnSe2 matrix, a relatively common occurrence in the sample. TEM results indicate that SnSe phases in the range from several to hundreds of nanometers exist in the SnSe2 matrix, and the thermoelectric properties of the material will be seriously influenced by the nano-composite structure of SnSe2/SnSe. In order to study the chemical state and shed light on the valence structure of Sn in SnSe2/SnSe nano-composites, four typical samples (undoped composite and 6% Cl-doped composite, undoped SnSe and 4% Cl-doped SnSe) were chosen for XPS measurement. Figure 5 shows the photoemission spectrum of the Sn-3d core state. As for the two composites, the shape of the peak is asymmetric indicating that Sn is in the mixed valence state in the structure. For the undoped SnSe2/SnSe composite, the binding energy of the large peak on the left which is drawn in black is 493.5 eV and it can be fitted by two smaller peaks with the binding energies of 493.4 eV (violet line) and 493.8 eV (dark blue line), which correspond to two core states of Sn4+ (3d3/2) and Sn2+ (3d3/2), respectively. The binding energy of the large peak on the right which is drawn in black is 485.25 eV and it can be fitted by two binding energies of smaller peaks of 485.10 eV (violet line) and 485.50 eV (dark blue line), which correspond to two core states of Sn4+ (3d5/2) and Sn2+ (3d5/2), respectively. As for the 6% Cl-doped composite, obviously the peak splitting here is significant and the peak can be viewed 12


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as consisting of four core states Sn4+(3d3/2), Sn2+(3d3/2), Sn4+(3d5/2), and Sn2+(3d5/2), demonstrating that all samples are SnSe2/SnSe composites that include SnSe2 and SnSe phases. After doping with Cl, the binding energy of the Sn2+ and Sn4+ core states shifts to the higher energy, indicating the presence of hybridization between Cl and Sn. The fact that the Cl-Sn binding energy is much larger than the Se-Sn binding energy further attests to Cl successfully entering into both SnSe2 and SnSe compounds. Nevertheless, the hybridization is more pronounced in the Sn4+ peak rather than that in the Sn2+ one, therefore, it is insensible to know whether Cl doped into the SnSe second phase and make it a p-n transition since it is well known that the SnSe phase can change from the p-type material to the n-type material due to doping with Cl. Based on this consideration, we also measured the X-ray Photoemission spectra of Cl doped SnSe and undoped SnSe without the SnSe2 phase. Indeed the peaks for Sn3d3/2 and Sn3d5/2 core states are symmetric without detection of Sn4+. Moreover, it should be noted that the binding energy of Sn2+ shifts toward higher energy upon doping with Cl and the shift of the bind energy for Sn2+ is not that obvious in comparison with that of the Sn4+ upon doping with Cl, which is similar to that we observe in the XPS of SnSe2/SnSe composite upon Cl doping, indicating Cl indeed doped into both the main phase of SnSe2 and the secondary phase of SnSe in the composites. Actually, the extent of the shift in the binding energy is a reflection of extent to what the hybridization between Sn and Cl, as a result, the intensity of hybridization between Sn2+ and Cl is not that strong as that between Sn4+ and Cl. 3.2 Thermoelectric performance 13



Structure and composition analyses show that Cl substitutes on the Se site. Since Cl has one more electron in its outermost shell than Se, it acts as a donor and increases the carrier concentration in the system. In the following sections we discuss the effect of Cl doping and the presence of nanoprecipitates on the transport properties of SnSe2-based composites. 3.2.1 Electronic properties at low temperatures Temperature dependence of the carrier concentration and the carrier mobility in the range of 10 K–300 K measured parallel to the pressing direction are shown in Figure 6. The carrier concentration of all samples is almost constant at low temperatures and it rises notably only above 220 K. The carrier concentration at 300 K increases with the increasing content of Cl doping from 4.3 × 1018 cm-3 for undoped SnSe2/SnSe composite to 5 × 1019 cm-3 for 6% Cl-doped composite, the latter value being well within the range of the optimum carrier concentration predicted by theoretical calculations.32 Using the measured carrier concentration and electrical conductivity, the calculated temperature dependence of the carrier mobility of all samples is shown in Figure 6(b). The undoped sample shows very weak temperature dependence in the whole measured temperature range, indicating the dominance of neutral impurity scattering. The Cl-doped samples display a modestly decreasing carrier mobility with temperature till about 220 K. At higher temperatures the mobility decreases more rapidly and follows an approximately T-3/2 dependence, suggesting that the carrier scattering becomes dominated by acoustic phonons. Surprisingly, upon Cl doping, the carrier mobility increases dramatically along both 14


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directions (parallel and perpendicular to the pressing direction applied during SPS). For example, the room temperature carrier mobility along the pressing direction rises from 2.6 cm-2V-1s-1 for the undoped sample to 9.8 cm-2V-1s-1 for sample with the nominal Cl content of 2.5. In the direction perpendicular to the pressing direction, the mobility increases from 7.4 cm-2V-1s-1 for undoped composite to 16.1 cm-2V-1s-1 for 1.5% Cl-doped composite. Hence, both the carrier concentration and the mobility are improved along both directions after doping with Cl, a somewhat unusual situation for a material. In a conventional semiconductor, the increased carrier concentration via doping is expected to enhance point defect scattering and potentially also carrier-carrier scattering, leading to a decrease in the carrier mobility. It should be noted that even if the structure possessed the orientation factor significantly above 0.5, this could not result in an enhancement of the carrier mobility in both parallel and perpendicular directions. In SnSe2/SnSe composites, heterogeneous interfaces formed by SnSe and SnSe2 are likely to affect the charge distribution and the electronic band alignment, which, in turn, would influence the charge transport. To shed more light on the issue, Ultraviolet Photoemission Spectroscopy (UPS) and Ultraviolet–Visible Spectroscopy (UV-VIS) were applied to x%Cl-SnSe2/SnSe composites (x=0, x=6). The results are shown in Figures 7(a) and 7(b). The work function (W) is obtained by the following formula: = ℎ − −

(2)

where hv is the incident photon energy. Since we used He I as the ionization source, the energy hv is 21.2 eV. The work function (W) refers to the minimum energy 15



required to remove an electron from a solid to the vacuum energy level.54 The work functions for undoped SnSe2/SnSe composite and 6% Cl-doped composite are 5.4 eV and 5.25 eV, respectively. The energy difference between the Fermi level and the highest occupied states (the valence band maximum) is 0.88 eV and 1.04 eV, respectively. UV-VIS measurements of the undoped composite and the 6% Cl-doped composite show that the electronic absorption edge shifts slightly to a lower energy after doping with Cl, from 1.05 eV for undoped composite to 1.0 eV for the 6% Cl-doped composite. Therefore, the configurations of the energy bands in undoped SnSe2/SnSe composite and in 6%Cl-SnSe2/SnSe composite are depicted in Figure 7(c). After doping with Cl, the Fermi level moves from the gap close to the conduction band minimum into the conduction band. The reduction in the work function W may be related to changes in SnSe2/SnSe heterojunctions. It is well known that intrinsic SnSe is a p-type semiconductor with the band gap of 0.86 eV,55 while intrinsic SnSe2 is an n-type semiconductor with the band gap of 1.05 eV. For undoped SnSe2/SnSe composites, an electron from n-type SnSe2 near the heterojunction diffuses to p-type SnSe, leaving behind a positively charged ion in n-type SnSe2 that recombines with a hole in p-type SnSe. Similarly, a hole in p-type SnSe diffuses to n-type SnSe2, leaving behind a negatively charged ion in p-type SnSe that recombines with an electron in n-type SnSe2. The region near the p–n interface loses its neutrality and most of its mobile carriers, giving rise to a space charge region or a depletion layer, as shown in Figure 7(e). Therefore, an electric field will be induced by the 16


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space charge region to oppose diffusion of both electrons and holes until the p-n heterojunction attains its thermal equilibrium. When electrons approach the p-n heterojunction, their trajectory and drift velocity will be hindered, and the band structure of the p-n interface will also be altered. Figure 7(d) presents a schematic band diagram of the heterogeneous interface of SnSe2/SnSe before and after Cl doping. SnSe2 is an n-type non-degenerate semiconductor with the Fermi level EFn close to the bottom of conduction band, while SnSe is a p-type non-degenerate semiconductor with the Fermi level EFp close to the top of the valence band. When the two compounds make a contact, electrons flow from the n-region with the Fermi level close to the conduction band minimum to the p-region with the Fermi level near the valence band maximum, and holes flow from the p-region to the n-region. As a result, EFn moves down and EFp moves up until the two Fermi levels equate to each other, EFn=EFp，giving rise to a uniform Fermi level EF in the p-n contact region, see Figure 7(d). Owing to band bending, electrons must overcome the potential difference ∆E1 of the Schottky junction as they move from the n-region to the p-region. Therefore, the presence of p-type SnSe will seriously affect the flow of electrons in the matrix. In the case of the Cl-doped SnSe2/SnSe composite, the SnSe phase changes from the p-type material to the n-type material due to doping with Cl. At the same time, SnSe2 becomes a heavily n-doped semiconductor. Thus, the Fermi level positions are altered and the nature of charge carrier conduction is different. Band bending, again, gives rise to the interfacial potential barrier ∆E2, but it is now much smaller than in the case of the undoped SnSe2/SnSe composite, as shown schematically in the lower panel of 17



Figure 7(d). The contact is now ohmic and the flow of charge carriers is considerably less hindered than in the former case, see a sketch in Figure 7(e). Consequently, the carrier mobility in Cl-doped composites is significantly enhanced. Table 1 collects room temperature parameters of x%Cl-SnSe2/SnSe composites, including the electrical conductivity and the carrier mobility measured along different directions. The numbers in bold refer to parameters measured in the direction parallel to the pressing direction, while the numbers in parenthesis refer to parameters measured in the direction perpendicular to the pressing direction. It is noteworthy that the room temperature electrical conductivity and carrier mobility measured on samples oriented perpendicular to the pressing direction is about a factor of two larger than for samples oriented parallel to the pressing direction, implying a quite high anisotropy of transport properties. Figure 8(a) depicts the electrical conductivity of composites measured on samples oriented parallel to the pressing direction. The undoped composite has weakly semiconducting temperature dependence while the Cl-doped composites show a distinctly decreasing electrical conductivity with the increasing temperature, the behavior characteristic of a highly degenerate semiconductor. As a result of the higher carrier concentration and improved carrier mobility, the room temperature electrical conductivity significantly increases from 200 Sm-1 for the undoped composite to 6.4×103 Sm-1 for 7% Cl-doped composite. All samples exhibit negative Seebeck coefficients over the entire temperature range and for both pressing directions (Figure 8(b)), documenting the n-type character of transport with electrons the dominant charge carrier. The absolute value of the 18


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Seebeck coefficient of the undoped composite and the two lowest doped composites increases with the rising temperature in the range 300 K - 650 K, reaches its maximum value, and then begins to decrease at the highest temperatures (650 K-773 K). The decrease in the value of the Seebeck coefficient is caused by the onset of intrinsic excitations at high enough temperatures. Composites doped with more Cl have high enough carrier concentration, which, in the temperature range covered, is not overwhelmed by the presence of minority carriers from intrinsic excitations. Consequently, their Seebeck coefficients do not show a turnaround at the highest temperatures. Of course, the magnitude of the Seebeck coefficient decreases with the increasing content of Cl on account of the progressively higher carrier concentration. Assuming that the average mean free path of charge carriers is constant, the Seebeck coefficient of degenerate semiconductors is described by the Mott equation: α=

∗ !

"#

$"%&

'( "

(3)

where α is the Seebeck coefficient, kB is the Boltzmann constant, h is the Planck constant, e is the electron charge, n is the carrier concentration, and m* is the effective mass. Although the single parabolic model might be an over-simplified picture, it should, nevertheless, adequately describe the trend in the transport properties of Cl-doped SnSe2/SnSe composites. Figure 8(c) shows Pisarenko plots of the measured Seebeck coefficient at 300 K of Cl-doped SnSe2/SnSe composites as a function of the carrier concentration. The curves are plots for different values of the effective mass and the experimental values of the Seebeck coefficient fall on the curve with the effective mass of 1.2 m0, where m0 is the free electron mass. Due to the enhanced 19



Page 20 of 46

electrical conductivity and the effective mass m*, the power factors of composites rise at 600 K with Cl doping from 0.07 mWm-1K-2 to 0.39 mWm-1K-2 along the pressing direction, as shown in Figure 8(d). 3.2.2 Thermal conductivity Figure 9(a) shows the temperature dependence of the total thermal conductivity for all composite samples along the pressing direction. As the temperature increases, the thermal conductivity of all samples decreases on account of the enhanced contribution of Umklapp processes. Moreover, the thermal conductivity decreases after Cl doping compared to the undoped SnSe2/SnSe composite. The thermal conductivity of heavily Cl-doped SnSe2/SnSe composite is very low in the whole temperature range due to the enhanced point defect scattering and heterogeneous interface scattering. The room temperature value is 1.39 Wm-1K-1 and it decreases down to 0.42 Wm-1K-1 at 773 K. We estimate the lattice thermal conductivity by subtracting the electronic thermal conductivity from the measured total thermal conductivity:

κ = κlat + κele

（ 4）

κele = LσT

（5）

Here, κlat is the lattice thermal conductivity, κele is the electronic thermal conductivity,

σ is the electrical conductivity, and L is the Lorenz number. Assuming a single parabolic band model and acoustic phonon scattering as being the dominant phonon scattering mechanism in the temperature range (the parameter r = -1/2), the Lorenz number L is calculated from:56 20


Page 21 of 46 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


' ,-⁄'

) = $ * & +,"⁄'/01 :=±

*

⁄ 23

/04⁄ 23

− 5 ' 67 8

, 67 = ?F

∞

8

@ A B@ ,CDE @23

(6) (7) (8)

where kB is the Boltzmann constant, e is the elemental electron charge, α is the Seebeck coefficient, and ηF is the reduced Fermi level. The obtained lattice thermal conductivity is shown in Figure 9(b). Comparing Figures 9(a) and 9(b), it follows that the lattice thermal conductivity measured parallel to the pressing direction is very close to the total thermal conductivity, indicating that the lattice thermal conductivity, by far, dominates the total thermal conductivity. After Cl doping, the lattice thermal conductivity has slightly decreased in comparison to the undoped composite, undoubtedly on account of enhanced point defect scattering. Combining the electrical and thermal transport properties, the ZT value was calculated and is shown in Figure 10. The figure of merit ZT increases with the increasing temperature. Doping with Cl has modified the heterojunction from the p-n type to the ohmic n-n type, significantly enhancing not just the carrier concentration but also the carrier mobility, resulting in a much improved power factor. Benefitting from the intrinsically low thermal conductivity of SnSe2 and SnSe and its further reduction upon Cl-doping, the 6%Cl-SnSe2/SnSe composite sample has the highest ZT value of 0.56 at 773 K in the direction parallel to the pressing direction. This is about 6 times higher than the ZT value of the undoped composite. Because both SnSe2 and SnSe are layered compounds, their physical properties are highly anisotropic and the influence of anisotropy on the thermoelectric properties 21



must be taken into account. Here we take the 6% Cl-doped composite as an example. As shown in Figure 11(a), the electrical conductivity along both the perpendicular and parallel pressing directions decreases with the increasing temperature. Notably, the electrical conductivity measured perpendicular to the pressing direction is much higher than the one measured parallel to the pressing direction. Specifically, at room temperature, the perpendicular electrical conductivity is 1.15×104 Sm-1, about twice the value of the parallel electrical conductivity (5×103 Sm-1), owing to a large difference in the carrier mobility in the two measured directions. Interestingly, there is very little anisotropy in the Seebeck coefficient, as documented in Figure 11(b). Due to the higher electrical conductivity in the direction perpendicular to the pressing direction, the power factor here is about an order of magnitude larger than in the direction parallel to the pressing direction, as shown in the inset of Figure 11(b). The maximum power factor of 0.75 mWm-1K-2 is attained at 600 K in the direction perpendicular to the pressing direction. Figure 11(c) shows the temperature dependence of the total and lattice thermal conductivities measured on samples oriented along the perpendicular direction with respect to the pressing direction, κ⊥ and κL,⊥, and along the parallel direction, κ// and κL,//. Both the total and lattice thermal conductivities in the direction perpendicular to the pressing direction are much larger than those measured in the parallel direction. For instance, at room temperature, the thermal conductivities κ⊥ and κL,⊥ are 3.8 Wm-1K-1 and 3.3 Wm-1K-1, respectively, while those in the parallel direction are κ// = 1.43 Wm-1K-1 and κL,// = 1.36 Wm-1K-1. The difference is largely attributable to a large inter-layer distance (~3.1 Å) 22


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forming the weak van der Waals bond, and thus a very low cross-layer sound velocity.57, 58 Although the electrical conductivity and the power factor are larger in the perpendicular direction, the much increased thermal conductivity in this direction leads to the lower ZT value of 0.38 at 773 K, as shown in Figure 11(d). Conclusions SnSe2/SnSe composites with uniformly distributed nano-sized SnSe were prepared by the melting-quenching process, and the influence of Cl doping on the thermoelectric properties of the nanocomposites was studied systematically. Doping with Cl modifies the charge distribution and decreases the energy barrier in the interfacial region between the two phases constituting the composite. Consequently, doping with Cl gives rise to a simultaneous increase in the carrier concentration and carrier mobility, significantly increasing the electrical conductivity. Furthermore, the enhanced point defect scattering on account of the presence of Cl atoms, together with effectively phonon scattering heterogeneous SnSe-SnSe2 interfaces, result in a very low room temperature lattice thermal conductivity of 1.43 Wm-1K-1. A composite doped with 6% of Cl reached the maximum ZT value of 0.56 at 773 K in the direction parallel to the pressing direction applied during the SPS process. This is the best result achieved with the SnSe2 system so far. Even though higher power factors can be realized with composites oriented perpendicular to the pressing direction, their thermal conductivity is, unfortunately, much higher than for samples oriented parallel to the pressing direction, and their ZT values are inferior in comparison. This is an unusual situation for the conventional two dimensional structure and contradictory to 23



theoretical estimates.32

ASSOCIATED CONTENT Supporting Information: The phase diagram of Sn-Se; The EDS spectrum analysis of the large white circular region in Figure 4(a); The electrical and thermal transport properties of SnSe1-xClx samples measured parallel to the pressing direction. Corresponding Author Email: [email protected] (X.Su), [email protected] (X.Tang) Notes The authors declare no competing financial interest.

ACKNOWLEDGMENTS

This work is financially supported by the Natural Science Foundation of China (Grant No. 51521001, and 51632006), the Fundamental Research Funds for the Central Universities (WUT: 162459002, 2015Ⅲ061) and the 111 Project of China (Grant No. B07040).

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Mater Today Phys 2017, 3, 127-136.

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Figure 1: (a) Crystal structure of SnSe2; (b) the layered structure of SnSe2 molecular

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Figure 2: (a) XRD patterns of all x%Cl-SnSe2/SnSe composites before SPS sintering; (b) expanded view of high-angle XRD patterns (59.6°-60.8°); (c) XRD patterns of bulk x%Cl-SnSe2/SnSe composite after SPS sintering; (d) orientation factors of all bulk x%Cl-SnSe2/SnSe composite samples perpendicular to the pressing direction along the (001) crystal surface.

35



Figure 3: (a) and (b) FESEM images of a freshly fractured surface of 6%Cl-SnSe2/SnSe composite taken at different magnifications; (c)-(f) Back Scattering Electron images of the polished surface and element maps collected by EDS for 6%Cl-SnSe2/SnSe composite.

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Figure 4: (a) A low-magnification TEM image of undoped sample showing different structures; inset of Figure (a) shows the selected area electron diffraction image of the square area; (b) an electron diffraction image of Figure (a) shows two kinds of different diffraction spots that refer to SnSe2 and SnSe, respectively; (c)-(d) HRTEM images of small circular areas in Figure (a) indicate the presence of SnSe nano-precipitates in the matrix of SnSe2.

37



Figure 5: Photoemission spectra of Sn3d3/2 and Sn3d5/2 core states of the samples .

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Figure 6: (a) Temperature dependence (10 K-300 K) of the carrier concentration; (b) Temperature dependence of carrier mobility of x%Cl-SnSe2/SnSe composite measured on samples oriented parallel to the pressing direction.

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Figure 7: (a) Ultraviolet photoelectron spectroscopy (UPS) for un-doped SnSe2/SnSe composite

and

6%Cl-SnSe2/SnSe

composite;

(b)

UV-VIS

spectra

of

x%Cl-SnSe2/SnSe composite (x=0, 6); (c) energy band structure of un-doped SnSe2/SnSe and 6%Cl-SnSe2/SnSe composites; (d) the schematic band diagram of homogeneous SnSe2/SnSe composite before and after Cl doping; (e) a sketch of the conduction behavior before and after Cl doping.

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Figure 8: Temperature dependence of (a) the electrical conductivity, (b) the Seebeck coefficient, and (c) carrier concentration dependence of the Seebeck coefficient at room temperature; (d) the power factor of x%Cl-SnSe2/SnSe composite measured on samples oriented parallel to the pressing direction.

41



Figure 9: Temperature dependence of (a) the total thermal conductivity, (b) the lattice thermal conductivity of x%Cl-SnSe2/SnSe composite measured on samples oriented parallel to the pressing direction.

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Figure 10: Temperature dependence of the thermoelectric figure of merit (ZT) of x%Cl-SnSe2/SnSe composite oriented parallel to the pressing direction.

43



Figure 11: Temperature dependence of (a) the electrical conductivity, (b) the Seebeck coefficient, (c) the thermal conductivity and (d) thermoelectric figure of merit (ZT) for the 6% Cl-doped composite measured parallel and perpendicular to the pressing direction. The power factor in the two respective directions is shown in the inset of Figure 11(b).

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Table 1: Room temperature parameters of x%Cl-SnSe2/SnSe composite (The numbers in bold refer to parameters measured in the direction parallel to the pressing direction. The numbers in parenthesis refer to parameters measured in the direction perpendicular to the pressing direction). µ α σ n L Sample F 18 -3 -2 -1 -1 3 -8 -2 x=0 x=1.5 x=2.5 x=3.5 x=5 x=6 x=7

(10 cm )

(cm V S )

4.3 28.8 34.9 39.8 41.7 45.9 50.8

2.8/(7.4) 8.8/(16.1) 9.8/(15.3) 7.7/(14.5) 7.2/(14.9) 7.2/(15.2) 7.1/(14.8)

(µV/K)

0.5 0.56 0.53 0.54 0.48 0.54 0.56

-415 -245 -216 -216 -204 -201 -204

45


(10 S/m)

0.2/(0.5) 4.1/(7.4) 5.5/(8.6) 4.9/(9.2) 4.9/(10.1) 5.3/(11.5) 6.4/(14.4)

(10 WK ) 1.49 1.56 1.59 1.59 1.61 1.61 1.61


Table of Contents graphic

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SnSe nano-com - ACS Publications

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