Enhancing Thermoelectric Performance of n-Type Hot Deformed

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Enhancing Thermoelectric Performance of n-type Hot Deformed Bismuth-Telluride-Based Solid Solutions by Non-stoichiometry Mediated Intrinsic Point Defects Renshuang Zhai, Lipeng Hu, Haijun Wu, Zhaojun Xu, Tie-Jun Zhu, and Xin-Bing Zhao ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b08537 • Publication Date (Web): 04 Aug 2017 Downloaded from http://pubs.acs.org on August 7, 2017

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ACS Applied Materials & Interfaces

Enhancing Thermoelectric Performance of n-type Hot Deformed Bismuth-Telluride-Based Solid Solutions by Non-stoichiometry Mediated Intrinsic Point Defects Zhai, Renshuang; † Hu, Lipeng; † Wu, Haijun; ‡,§ Xu, Zhaojun; † Zhu, Tie-Jun; †,* Zhao, Xin-Bing † †

State Key Laboratory of Silicon Materials, School of Materials Science and Engineering,

Zhejiang University, Hangzhou 310027, China ‡

Department of Physics and Shenzhen Key Laboratory of Thermoelectric Materials, South

University of Science and Technology of China, Shenzhen 518055, China §

Department of Materials Science and Engineering, National University of Singapore, 7

Engineering Drive 1, Singapore 117575, Singapore * Correspondence: Professor T. J. Zhu, E-mail: [email protected]

KEYWORDS: Thermoelectric materials, non-stoichiometry, bismuth telluride, hot deformation, point defect, Se deficiency.

ACS Paragon Plus Environment

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ABSTRACT Bismuth-telluride-based solid solutions are the unique thermoelectric materials near room temperature. Various approaches have been applied to enhance the thermoelectric performance and much progress has been made in their p-type materials. However, for the n-type counterparts, little breakthrough has been obtained. We herein report on enhancing thermoelectric performance of n-type bismuth-telluride-based alloys by non-stoichiometry to mediate the point defects, combined with one-time hot deformation. The improved power factor of 3.3×10-3 Wm-1K-2 and reduced lattice thermal conductivity contribute to a high figure of merit zT of 1.2 at 450 K for n-type Bi2Te2.3Se0.69 alloys with Se deficiency. The high zT is comparable to that of Bi2Te2.3Se0.7 hot deformed for three times, which is a practically complicated process. The results demonstrate that non-stoichiometry can be an effective and simple strategy in mediating intrinsic point defects and enhancing the thermoelectric performance of bismuthtelluride-based alloys.

INTRODUCTION Thermoelectric (TE) devices have drawn extensive attention over the past decades due to the potential in the direct inter-conversion between thermal and electrical energy. The conversion efficiency depends on the dimensionless figure of merit zT of the TE materials, zT = α2σT/κ, where α is the Seebeck coefficient, σ and κ are the electrical conductivity and thermal conductivity (including the carrier contribution κel and the lattice contribution κph) respectively, and T is the operating temperature. The TE parameters α, σ, and κel are interrelated closely via carrier concentration n, and κph is considered to be a relatively independent parameter. Phonon engineering1-4 and band engineering4-7 have been proved two main strategies to enhance zT,


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aiming to decrease lattice thermal conductivity κph and boost the electrical power factor (PF = α2σ), respectively. Bismuth-telluride-based alloys are the best commercially available TE materials for solidstate refrigeration and power generation near room temperature. These highly orientated bismuth-telluride-based ingots are generally grown by zone melting (ZM) method. The unidirectionally solidified ingots with strong texture exhibit an extraordinary PF,8 leading to a high in-plane zT around unity near room temperature. However, the relatively high in-plane intrinsic κph limits the zT of ZM ingots. In addition, the poor mechanical property of ZM ingots erects a barrier to device manufacturing.9 For p-type bismuth-telluride-based alloys, nanostructure engineering has been proved effective in enhancing TE performance via grain refinement and various nano-defects, which suppress κph pronouncedly with less detrimental influence on electrical transport. The “bottomup” nanostructuring approach refers to various synthesis routes of nano-scale precursors such as ball milling,1 hydrothermal method,10 melt spinning,11 microwave-stimulated wet-chemical method,12 etc., followed by the sintering methods of hot pressing (HP) or spark plasma sintering (SPS). Through a “bottom-up” approach, the maximum zT of p-type bismuth-telluride-based alloys has reached > 1.4 around 400 K. In recent years, the “top-down” approach has been developed by which the nanostructures are in situ formed in the bulk TE materials. Shen et al. reported that in situ nanostructures can form in the hot deformed Bi0.5Sb1.5Te3 alloys with a peak zT > 1.3 near room temperature.13 Zhu et al. directly hot deformed the p-type Bi2-xSbxTe3 ZM ingots, resulting in a high average zT ~ 1.2 in the temperature range of 300 K ~ 525 K.14, 15 Subsequently, they investigated the application of hot deformation (HD) to bismuth-telluridebased alloys for low- and mid-temperature power generation, such as hot deformed p-type Bi2-


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16 xSbxTe3,

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p-type In-alloyed Sb2Te317 and n-type Bi2Te2Se18. Recently, Xu et al. attained high

performance p-type bismuth-telluride-based alloys with a peak zT ~ 1.4 at 500 K and an average zT ~ 1.3 in the temperature range of 400 K ~ 600 K via direct hot deformation of In-alloyed Bi0.3Sb1.7Te3 ZM ingots, due to the synergistic role of point defects engineering, band structure engineering and multiscale microstructuring.19 On the basis of nanostructure engineering, dislocation engineering2 and dispersed nano-particles addition20 were also introduced to scatter phonons. Remarkable progress has been made in p-type bismuth-telluride-based polycrystalline alloys in the last decades. However, nanostructure engineering is not favourable for the n-type bismuth-telluride-based alloys due to the loss of texture. The intrinsic electrical anisotropy of n-type bismuth-telluridebased alloys is stronger than p-type counterparts21, leading to a larger deterioration in PF and consequently impaired zT. Xie et al. combined melt spinning (MS) with SPS to obtain a zT ~ 1.05 at 420 K.22 Wang et al. reported Te-free n-type bismuth-telluride-based alloys with a zT ~ 1 at 800 K, via phonon softening and band converging.23 Much work has been focused on the texturing engineering for n-type bismuth-telluride-based alloys to maintain carrier mobility. Zhao et al. utilized the two-step SPS to enhance the texture and the obtained alloy exhibited an increment of about 20% in zT compared to the directly SPS counterparts.24 Yan et al. re-pressed the hot-pressed Bi2Te2.7Se0.3 to improve zT from 0.85 to 1.04 at 400 K.25 Liu et al. re-pressed the Cu0.01Bi2Te2.7Se0.3 HP ingot to obtain a zT ~ 1.06 at around 400 K.26 Pan et al. utilized the synergistic effect of texturing and nanostructuring to achieve a zT ~ 1.1 at 473 K.27 Hu et al. demonstrated that repetitive hot deformations enhanced texture pronouncedly.28 Tang et al. conducted three times hot deformations on 0.1 at.% SbI3–Bi2Te1.9Se1.1 alloy, which exhibits a zT ~ 1.1 at about 600 K.29 Based on the HD process, Hu et al. manipulated intrinsic point defect to


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tailor carrier concentration and induced multi-scale microstructure to scatter phonons. The texture was enhanced due to HD three-times, reaching a state-of-the-art peak zT ~ 1.2 at 445 K.30 Nevertheless, the complicated procedure of multi-HD limits the application. Carrier concentration can be adjusted by hot deformation, due to the deformation-induced donor-like effect and the thermal recovery role.30,

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Non-stoichiometry is involved in the

formation of intrinsic point defects. According to the first principles calculation32, the growth condition of rich cations makes Sb′Te , Bi′Te and VSe•• dominant point defects in Sb2Te3, Bi2Te3 and Bi2Se3 respectively; On the contrast, VSb′′′ , Te•Bi , Se•Bi are dominant accordingly under an anion-rich condition. Horak et al. reported that cation antisite defect Bi′Te dominated in Bi2Te3 single crystal prepared by Bridgeman method conducting as p-type semiconductor with hole concentration around 6-7×1018 cm-3, while anion vacancy VSe•• played a dominant role in the single crystal Bi2Se3 prepared by the same method and the alloy exhibits n-type behavior with electron concentration about 1-2×1019 cm-3.33 Excess Bi in Bi2Te3 single crystal results in the formation energy decrease and more antisite defects Bi′Te and hence higher hole concentration. Similarly, excess Bi in Bi2Se3 induced more antiste defects Bi′Se into lattice and decreased the electron concentration.34 In addition, the recent investigation on topological insulator has proved the dependence of carrier concentration on non-stoichiometry. Bi2Te2Se1 single crystal, in the terms of composition, is in the vicinity of the transformation point from p-type to n-type behavior, making a good basic matrix for topological insulator.35 However, it always exhibits ntype behavior due to the vacancies of anions in Bi2Te2Se1 prepared according to the nominal composition. Jia et al. prepared high-resistivity Bi2+xTe2-xSe1 samples by inducing extra Bi′Te


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into lattice to lower electron concentration.36 However, utilizing non-stoichiometry to adjust point defects has rarely been reported in this thermoelectric polycrystalline alloys. In this vein, non-stoichiometry was employed to mediate intrinsic point defects of n-type bismuth-telluride-based alloys in this work, together with HD processing, aiming to further enhancing the TE performance. Only one-time HD was performed on Bi2Te2.3Se0.7-x (x = 0-0.05) samples. The donor-like effect, the enhanced texture and the simultaneously decreased κph all contribute to a high zT ~ 1.2 at 450 K in Bi2Te2.3Se0.69, highlighting the non-stoichiometry as an effective strategy in boosting thermoelectric figure of merit of hot deformed bismuth-telluridebased alloys.

EXPERIMENT Highly pure element chunks of Bi (5N), Te (5N), Se (5N) were weighed according to the nominal composition Bi2Te2.3Se0.7-x (x = 0, 0.01, 0.02, 0.03, 0.04, 0.05) and sealed into wellcleaned quartz tube at 10-3 Pa. The mixture was sequentially melted in a Muffle furnace. The melts were rocked every several hours to make sure of the homogeneity of the initial Bi2Te2.3Se0.7- x ingots. Fine powders were obtained by ball milling the ingot and sieved to 300mesh. The powders were hot pressed into a cylinder of Φ 12.7 mm in the graphite die at 673 K for 30 min, and the bulk samples are named as HP. Subsequently, a hot deformation process was performed by repressing the bulk into a disk of Φ 20 mm at 823 K for 30 min (called as HD samples). The typical photos of the bulk samples before and after hot deformation can be referred to our previous work.13 More details can be found elsewhere.28 Phase structure of all the samples were examined by X-ray Diffraction (XRD) with a Rigaku D/MAX-2550P diffractometer. Transmission Electron Microscope (TEM) observation


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was carried out on JEOL 2100F and FEI TF30 microscopes. The thin TEM specimens were prepared by the conventional standard methods including cutting, grinding, dimpling, polishing, and Ar-ion milling on a liquid nitrogen cooling stage. Netzsch LFA 457 laser flash apparatus with a Pyroceram standard was used to measure the thermal diffusivity (D). The specific heat (CP) was measured on the Netzsch DSC 404C and the density (ρ) was estimated by an ordinary dimension-and-weight measurement procedure. The thermal conductivity was then computed according to the relation κ = DρCP. The in-plane electrical conductivity (σ) and the Seebeek coefficient (α) were simultaneously measured on a commercial Linseis LSR-3 system. The Hall coefficient (RH) was determined at 300 K on a Quantum Design PPMS-9T instrument using a four-probe configuration. Then the Hall carrier concentration (nH) and in-plane Hall mobility (µH) were calculated according to nH = 1/eRH and

µH = σRH, respectively. Note that the thermal and electrical properties of all the HP and HD samples were measured along the in-plane direction.

RESULTS AND DISCUSSION Point defects and carrier concentration No trace of impurity phases was found in all the samples as indicated in XRD patterns (Figure S1). Figure 1 displays the variation of carrier concentration (n) with Se deficiency in polycrystalline samples. Obviously, all the samples exhibit n-type conduction behavior. The monotonicity of the decrease of n in HP polycrystalline samples can be elucidated via point defect scenario. Two main processes are associated with the variations of point defect type and concentration in the deformed bulks: one is the non-basal plane slip created by mechanical deformation, inducing the vacancies of 2 VBi′′′ : 3 VSe•• 37; On the other hand, the accumulated VBi′′′


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around Bi′Se drives Bi atoms easily to diffuse back to initial positions, resulting in excessive VSe•• ,38 which is depictured as the donor-like effect: 2VBi′′′ +3VSe•• +Bi′Se → VBi′′′ +Bi×Bi +4VSe•• +6e′ .39 The donor-like effect mainly depends on the concentration of Bi′Se , which is formed during melting, considering that the same deformation condition in this work retains the concentration of VBi′′′ and VSe•• induced via mechanical deformation. The concentration of Bi′Se increases in the Bi2Te2.3Se0.7-x ingot with increasing the deficiency of Se. At x = 0, low concentration of antisite defects was depleted during ball milling. While at x > 0, the concentration of antisite defects increases and Bi′Se partially involves in the donor-like effect during hot pressing. Extra holes induced by antisite defects compensate for partial electrons induced via donor-like effect, resulting in the decrease of electron concentration n. When x increases from 0 to 0.05, n decreases from 7.2×1019 cm-3 to 1.3×1019 cm-3. Figure 1 also displays that n increases first and then decreases with increasing Se deficiency x in HD samples (discussed via following Figure 2), and the dependence of n on Se deficiency is repeatable according to our systematic work (Figure

S2).

Figure 1. Carrier concentration as a function of Se deficiency in the hot pressed and hot deformed Bi2Te2.3Se0.7-x polycrystalline alloys at room temperature.


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Figure 2 shows the point defects variation in Bi2Te2.3Se0.7-x ingots with the alloy compositions and processing histories. At x = 0, the concentration of antisite defects in the samples is relatively low, and the donor-like effect induced by ball milling consumes Bi′Se up. Recovery role is dominant during subsequent hot deformation process where the donor-like effect cannot be provoked and n decreases after HD, as Figure 1 displays. The concentration of antisite defects is relatively high in the ingot x = 0.01. The donor-like effect involves just part of antisite defects during ball milling, which is provoked further in the subsequent hot deformation process, inducing more extra electrons. When 0.02 < x < 0.04, sufficiently high concentration of antisite defects remains even after hot deformation, where the donor-like effect is provoked. Remaining antisite defects compensate for partial donor-like defects, contributing to the electron concentration decreasing along with increasing the deficiency of Se (0.02 < x < 0.04). At x > •• 0.04, the concentration of antisite defects is high enough to nearly deplete VBi′′′ and VTe induced

during ball milling. All similar phenomena were reported in Bi2-xSbxTe3 and Bi2Te3-xSex.30 It should be mentioned that Bi-rich growth environment is beneficial to form septuple-layer structure defects Bi3Te 4′ . Horak supposed that Bi′Te formed with slightly excessive Bi, while Bi3Te 4′ formed on the condition of sufficiently excessive Bi.34 The increase of Bi3Se 4′ and the decrease of Bi′Se with increasing Se deficiency may be another explanation for the variation of carrier concentration between HP and HD samples, considering Bi3Se 4′ defects make no contribution to the donor-like effect.


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Figure 2. Schematic of point defect variation of Bi2Te2.3Se0.7-x ingots with the alloy compositions and processing histories.

Electrical transport properties Improved electrical conductivity was obtained via non-stoichiometry (x = 0.01 and 0.02) to manipulate point defects, as shown in Figure 3. The variation of electrical conductivity is consistent with that of carrier concentration as a function of Se shortage. The Bi2Te2.3Se0.69 (x = 0.01) HD sample reaches the electrical conductivity peak of around 1.45×105 Sm-1 at room temperature, which is higher than that of Bi2Te2.3Se0.7 sample HDed three times (Figure 3(a)).

Figure 3(b) shows the electrical conductivities of the HD samples are all higher than the HP counterparts at different contents of Se deficiency. The typical in-plane XRD pattern in Figure

4(a) shows hot deformation enhances the (00l) texture,24 which improves the in-plane carrier mobility in polycrystalline Bi2Te2.3Se0.7-x, as in Figure 4(b). Particularly, the carrier


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concentration of Bi2Te2.3Se0.69 (x = 0.01) is boosted significantly due to the donor-like effect during hot deformation (Figure 1), contributing to an increase by ~70% of electrical conductivity.

Figure 5(a) displays the temperature dependence of α for Bi2Te2.3Se0.7-x samples. The Seebeck coefficient of all samples is negative, exhibiting n-type conduction behavior. When x increases from 0 to 0.02, α remains almost the same due to only slight variation in carrier concentration. However, α decreases with x > 0.02 due to the sharp decrease of n. Particularly, α of Bi2Te2.3Se0.69 (x = 0.01) HD sample is slightly smaller than that of Bi2Te2.3Se0.7 HDed three times. While x > 0.04, α sharply increases due to the decrease of n originating from compensation role during hot deformation.

Figure 3. (a) Temperature dependence of in-plane electrical conductivity of the hot deformed Bi2Te2.3Se0.7-x samples. (b) Se deficiency dependence of the room temperature electrical conductivity of the HP and HD Bi2Te2.3Se0.7-x samples.


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Figure 4. (a) Typical X-ray diffraction patterns of the Bi2Te2.3Se0.67 bulk samples before and after hot deformation. (b) Se deficiency dependence of the in-plane carrier mobility at room temperature of the HP and HD Bi2Te2.3Se0.7-x samples.

Figure 5. Temperature dependences of (a) Seebeck coefficient and (b) power factor of the Bi2Te2.3Se0.7-x HD samples.


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Figure 5(b) presents the temperature dependence of power factor (PF) for the HD polycrystalline samples. Consistent with electrical conductivity, PF increases first and then decreases when x > 0.03 with increasing the Se deficiency. The PF of HD Bi2Te2.3Se0.69 reaches 3.3×10-3 Wm-1K-2 at room temperature, comparable to Bi2Te2.3Se0.7 hot deformed three times. The high PF demonstrates non-stoichiometry is a feasible and simplified approach to achieve improved electrical properties comparable to those by complicated multiple hot deformations.

Thermal conductivity and microstructure Non-stoichiometry also has a significant influence on thermal conductivity. Figure 6(a) displays the temperature dependence of total thermal conductivity. With increasing x, κ at room temperature initially remains 1.3 Wm-1K-1 (0 < x < 0.02), subsequently decreases mainly due to the deterioration of the contribution of electron thermal conductivity (x > 0.02). The total κ of Bi2Te2.3Se0.69 HD sample is slightly higher than Bi2Te2.3Se0.7 sample hot deformed three times. Further, κ decreases initially and a climb emerges on account of intrinsic transport with increasing temperature. With deficiency of Se, the inflexion point of κ shifts to lower temperature in accordance with the variation of α. The minority carrier concentration declines, enhancing the detrimental influence of bipolar diffusion.

Figure 6(b) exhibits the variation of κph as a function of temperature calculated according to κph = κ-κel, where κel = LσT with L = 2×10-8 V2K-2. Room temperature κph of Bi2Te2.3Se0.7-x HD samples decreases to 0.40 Wm-1K-1 (x = 0.02), then climbs up with increasing x. Two possible explanations associated with point defects are discussed: one is the interaction of ansite defects and vacancies induced via mechanical deformation; On the other hand, sufficiently Bi-rich condition is preferable for the formation of line stacking-fault defects with Bi3Se 4′ septuple


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layers. From the former view, deficiency of Se induces considerable Bi′Se antisite defects into lattice, enhancing the scattering of phonons and lowering κph pronouncedly.40 It is noteworthy that κph of Bi2Te2.3Se0.7-x HD samples displays a relatively notable decrease at x = 0.01-0.02, due to the donor-like defects induced during HD process. Dominant recovery role eliminates partial donor-like defects and lattice line defects, exhibiting a moderately reduced κph when x > 0.03 or x = 0.

Figure 6. Temperature dependences of (a) thermal conductivity κ and (b) lattice thermal conductivity κph of Bi2Te2.3Se0.7-x HD samples. To confirm the possibility that Bi3Se 4′ septuple layers formed thus affected the transport of electron and phonons, and figure out the structural origin of both excellent carrier concentration and lattice thermal conductivity of Bi2Te2.3Se0.7-x (x = 0.01 or 0.02) compared with highly Sedeficient ones (e.g. x = 0.05), we investigated the microstructure of Bi2Te2.3Se0.7-x (x = 0.02 and 0.05) HD sample via TEM. Multiscale microstructural defects have been observed in both the samples including nano-scale strain-field domains (Figure S3), submicro-scale twin boundaries


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(Figure 7), and micro-scale grain boundaries (Figure S4), and similar observations were reported previously.13, 27, 41 The twins and twin boundaries are shown in Figure 7. Generally, twins forms to relax the high strain energy,39 in the present case, due to deformation. The density of twins of the present one-time deformed materials as shown in Figure 7(a) are relatively higher than those of pristine ingots and also heavily three-time deformed ones (instead, there are high density of high-strained distortions).41 Electron diffraction pattern along [110] zone axis from the twin boundary clearly reflect the crystallographic relations between the twin variants: the unsplitted reflections 00l (l=3n, n is integer) is corresponding to the twin plane; the splitted reflections along [00l] direction are for the deviated planes from two twin variants. The intensity ratio between the splitted spots in Figure 7(b) can be consistent the relative area ratio of two twin variants as shown in Figure 7(a). The mediate-magnification HRTEM image and its FFT image focus on one twin boundary, Figure 7(c, d), reflecting the similar information as low-magnification but can do more deep data analysis. Separately selecting the individual reflections (e.g., -11-5) from different twin variants can be Fourier transformed into its respective real images as the IFFT images shows in Figure 7(e, f). While the IFFT image from the unsplitted reflections can highlight the twin plane, Figure 7(g). The HRTEM lattice image in Figure 7(h, i) can clearly show the two twin variants and their twin boundary in atomic-scale. Twin boundaries can act as an alternative phonon scattering source while having a negligible effect on the electronic transport since the twin boundary is always coherent interfaces, which cannot disrupt the electronic transport with very short wavelength.42 For Bi2Te2.3Se0.7-x (x = 0.05) sample, some stacking faults can be found, as shown in Figure

8(a, b). The concentration of the stacking faults is not high, thus it is difficult to differentiate it in


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the electron diffraction pattern, the insets in Figure 8(a, b). HRTEM was employed to see the details inside the stacking fault, Figure 8(c). Actually, the stacking fault is one set of septuple layers within the quintuple-layer matrix lattices. The fast Fourier transformed image (FFT) of

Figure 8(c) can reflect locally reciprocal information of such defects; diffused reflection spots can be found between the normal 1/5 superlattices along [001] direction, Figure 8(d). To reflect the corresponding relationship between septuple layers and the diffused reflection spots, the inverse FFT (IFFT) images from 006 reflection spots, the diffused lines between 006 and 009 reflection spots, and 009 reflection spots can be shown in Figure 8(e-g). It is clear that the diffused lines are from the septuple-layer stacking faults, and no obvious ones were observed at x = 0.02 (Figure S5) in this work. With increasing the deficiency of Se, large amounts of Bi′Se antisite point defects would prefer to gather together to form line stacking-fault defects with Bi3Se 4′ septuple layers. The decrease in the concentration of Bi′Se antisite point defects will weaken donor-like effect and also the phonon scattering, which is one possible reason why heavy Se deficiency (x = 0.05) does not exhibit a higher carrier concentration and a lower thermal conductivity compared with the modest one (x = 0.02).


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Figure 7. Structural characterization of twin boundaries: (a) low-magnification TEM image with several twins; (b) electron diffraction pattern from a twin boundary, two sets of patterns are marked; (c) mediate-magnification image of a twin boundary; (d) FFT image of (c). (e, f) IFFT image of (c) from two sets of splitted -11-5 reflections; (g) IFFT image of (c) from the unsplitted 003 reflections; (h) HRTEM lattice image of a twin boundary; (i) enlarged image from the region in (h).


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Figure 8. Structural characterization of Bi2Te2.3Se0.7-x (x = 0.05) sample: (a) low-magnification TEM image showing stacking faults, marked with white dashed squares; (b) mediatemagnification TEM image focuses on one stacking fault, the region marked with a yellow circle is high-strain state and the inset electron diffraction pattern; (c) HRTEM lattice image showing one septuple layer within quintuple-layer lattices; (d) FFT image of (c) showing addition diffused reflection spots along [001], between normal 1/5 superlattices; (e-g) IFFT image from 006 reflection spots, the diffused lines between 006 and 009 reflection spots, and 009 reflection spots.

Dimensionless figure-of-merit zT Figure 9 displays the temperature dependence of dimensionless figure of merit zT of Bi2Te2.3Se0.7-x HD samples. Slight deficiency of Se lowers κph and meanwhile improves electrical


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conductivity, enhancing zT, due to optimized concentration of Bi′Se and donor-like effect via hot deformation. On the contrast, excessive deficiency of Se induces superfluous Bi′Se antisite defects, which compensate for donor-like defects, resulting in low n and eventually a relatively poor TE property. Moreover, excessively low n intensifies the detrimental influence of intrinsic conducting on TE property. Through non-stoichiometry combined with one-time hot deformation, the peak zT of Bi2Te2.3Se0.69 HD sample reaches 1.2 at 450 K, which exhibits a 34% increment compared with Bi2Te2.3Se0.7 HD sample. The state-of-the-art zT ~ 1.2 is comparable to that of the three-time hot deformed (HD3), and higher than those of the bulk polycrystalline counterparts prepared via various other techniques, as shown in Figure 10.

Figure 9. Temperature dependence of zT values of Bi2Te2.3Se0.7-x HD samples.


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Figure 10. Temperature dependence of zT values of n-type Bi2Te3-based alloys prepared via various techniques 22, 23, 25, 27, 30, 41, 43 (ZM: zone melting; HD: hot deformation; HD3: three-time hot deformations; HP: hot pressing; BM: ball milling; MS: melt spinning; SPS: spark plasma sintering).

CONCLUSIONS The non-stoichiometry is effective to manipulate point defects in polycrystalline n-type Bi2Te2.3Se0.7. Combined with one-time HD in Bi2Te2.3Se0.69, the carrier concentration increases due to the provoked donor-like effect, meanwhile the texture is enhanced and an improved PF is obtained. Proper Se deficiency (x = 0.01) induces sufficient point defects to scatter phonons effectively and reduced κph is obtained. The integrated effect in the n-type Bi2Te2.3Se0.69 (x = 0.01) HD polycrystalline alloy results in high zT ~ 1.2 at 450 K, about 34% increment over the pristine Bi2Te2.3Se0.7 HD sample, which is comparable to Bi2Te2.3Se0.7 sample HDed for three times. The present results highlight the proper non-stoichiometry to manipulate point defects can improve


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electrical transport properties and simplify the hot deformation processing. Besides, the strategy may shed light on the possibility of manipulating point defects in other TE materials.

SUPPORTING IMFORMATION XRD patterns of bulk HP-series and HD-series samples (Figure S1), Dependences of carrier concentration on the systematic non-stoichiometry of Bi2Te2.3Se0.7 (Figure S2), Structural characterization of nanoscale strain filed domain in typical HD samples (Figure S3), Structural characterization of grain boundaries in typical HD samples (Figure S4), Structural characterization of Bi2Te2.3Se0.7-x (x = 0.02) sample (Figure S5).

ACKNOWLEDGEMENTS The authors would like to thank Prof. Jiaqing He, from South University of Science and Technology of China, for the helpful discussion on the TEM observation. This work was supported by the Natural Science Foundation of China (61534001, 11574267 and 51571177).

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TABLE OF CONTENTS


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Enhancing Thermoelectric Performance of n-Type Hot Deformed

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