Shifting up the optimum figure of merit of p -type bismuth telluride-based thermoelectric materials for power generation by suppressing intrinsic conduction

The abundance of low-temperature waste heat produced by industry and automobile exhaust necessitates the development of power generation with thermoelectric (TE) materials. Commercially available bismuth telluride-based alloys are generally used near room temperature. Materials that are composed of p-type bismuth telluride, which are suitable for low-temperature power generation (near 380 K), were successfully obtained through Sb-alloying, which suppresses detrimental intrinsic conduction at elevated temperatures by increasing hole concentrations and material band gaps. Furthermore, hot deformation (HD)-induced multi-scale microstructures were successfully realized in the high-performance p-type TE materials. Enhanced textures and donor-like effects all contributed to improved electrical transport properties. Multiple phonon scattering centers, including local nanostructures induced by dynamic recrystallization and high-density lattice defects, significantly reduced the lattice thermal conductivity. These combined effects resulted in observable improvement of ZT over the entire temperature range, with all TE parameters measured along the in-plane direction. The maximum ZT of 1.3 for the hot-deformed Bi0.3Sb1.7Te3 alloy was reached at 380 K, whereas the average ZTav of 1.18 was found in the range of 300–480 K, indicating potential for application in low-temperature TE power generation. Thermoelectric materials, which convert temperature differences and electric voltage into each other, serve in refrigeration or power generation applications. Currently, bismuth telluride (Bi2Te3) and its alloys are the most widely used thermoelectric materials. Tie-Jun Zhu, Xin-Bing Zhao and co-workers from Zhejiang University, China, have now investigated the effect of antimony (Sb) alloying on bismuth tellurides through a series of polycrystalline solid solutions of Bi2-xSbxTe3—where x varies between 1.4 and 1.8—prepared by hot deformation. Systematic tuning of the alloy composition showed that higher antimony content raised the material's optimal conversion temperature by repressing undesirable conduction. This effect arises from an increase in both the hole concentration and the band gap in the material. For a composition where x is 1.7, the alloy showed optimal performances at 380 kelvin—a suitable temperature for low-temperature power generation from the waste heat generated by industry or vehicles. The p-type bismuth telluride-based polycrystalline materials suiting for low-temperature power generations (near 380 K) have been obtained through Sb-alloying and HD, which suppresses the detrimental effect of intrinsic conduction at elevated temperature via increasing the hole concentration and band gap. The hot-deformed Bi0.3Sb1.7Te3 alloy, not usual composition Bi0.5Sb1.5Te3, shows a maximum ZT of 1.3 at 380 K, indicating a bright application potential in low-temperature power generations.

INTRODUCTION Thermoelectric (TE) devices have attracted extensive interest over the past few decades because of their potential use in direct thermal-toelectrical energy conversion and solid-state refrigeration. The TE conversion efficiency of a material can be gauged by the dimensionless figure of merit ZT ¼ a 2 sT/k, where a, s, k and T are the Seebeck coefficient, the electrical conductivity, the thermal conductivity and the operating temperature, respectively. 1 Continuous effort has been invested toward improving the ZT values of TE materials, resulting in significant advances through phonon engineering 2-9 and band engineering. [10][11][12][13][14] For example, remarkable increases in ZT have been achieved in bulk nanomaterials via the enhancement of phonon scattering at boundaries to reduce lattice thermal conductivities. 2,4,6,7 Currently, the best commercial TE materials near room temperature are still rhombohedral bismuth tellurides and related solid solutions fabricated by unidirectional crystal growth. [15][16][17] Nanostructuring strategies have been devised to prepare highperformance bismuth telluride-based alloys, including bottom-up and top-down approaches. In the bottom-up approach, nanostructures are first obtained by low-temperature hydrothermal synthesis, 2,5 ball milling 6,18,19 or melt spinning 4,20 and are subsequently sintered and consolidated by hot pressing (HP) or spark plasma sintering to yield bulk alloys. Recently, Shen et al. 7,21 have developed a novel top-down approach to improve the TE performance of p-type bismuth telluride-based alloys utilizing hot deformation (HD)-induced in situ nanostructures and high-density lattice defects. More recently, we found that HD can also be used to improve the TE performance of n-type bismuth telluride-based materials because of multi-scale microstructural effects including microscale texture enhancement, lattice defects and donor-like effects at the atomic scale, and recrystallization-induced local nanostructures. 22,23 The ternary alloy Bi 0.5 Sb 1. 5 Te 3 has been so far the bestknown p-type materials for room-temperature refrigeration. Various advanced powder metallurgical techniques 4,5,7 have been implemented to induce high ZT values (approximately 1-1.3) in polycrystalline Bi 0.5 Sb 1.5 Te 3 alloys through structural modification.
However, maximum ZT values are only obtained at room temperature, as rising temperatures dramatically decreased ZT because of intrinsic excitation. 4,7,24 For low-temperature power generation applications below 500 K, the maximum ZT of p-type bismuth telluride-based alloys must be shifted to higher temperatures, employing a composition other than Bi 0.5 Sb 1.5 Te 3 . One feasible solution toward the optimization of ZT at elevated temperatures, making them suitable for low-temperature power generation, is to actively suppress intrinsic conduction at elevated temperatures via two strategies. The first is to broaden the band gap; the second is to increase the hole concentration. The origin of hole concentration in p-type Bi 2 Àx Sb x Te 3 mixed crystals is mainly associated with Sb 0 Te and Bi 0 Te antisite defects, [25][26][27] which are created when cations occupy the vacant anion sites during crystal growth from a stoichiometric melt. The formation energy of an antisite defect mainly depends on bond polarity. The lower the polarity of the bonds between cations and anions, the lower the energy for antisite defects. 26,27 Raising the Sb content in Bi 2 Àx Sb x Te 3 mixed crystals can reduce the formation energy of antisite defects, thereby increasing the hole concentration because of the smaller difference in electronegativity for Sb-Te compared with Bi-Te. 27 The hole concentration may be further tuned through donor-like effects induced during the deformation process by grinding and pressing. 22,23,28,29 The binary compound Bi 2 Te 3 has a narrow band gap (approximately 0.13 eV), [30][31][32][33] which can be adjusted through the formation of solid solutions by alloying with Sb 2 Te 3 (a larger band gap of approximately 0.28 eV 34,35 ). 36 The band gap increases with the addition of Sb to the rich-Sb 2 Te 3 region, which can be attributed to increasing degeneracy of charge carriers and the larger band gap of Sb 2 Te 3 compared with Bi 2 Te 3 . 35 Thus, Sb substitution can effectively shift the optimum ZT to higher temperatures. 37 Nonetheless, systematic work on the effects of Sb-substitution on the TE performance of polycrystalline Bi 2 Àx Sb x Te 3 is rare.
In this work, we fabricated a series of p-type Bi 2 Àx Sb x Te 3 (x ¼ 1.4-1.8) polycrystalline solid solutions by a top-down, HD method to explore the effect of Sb-alloying on the suppression of intrinsic conduction. The optimum ZT was shifted to higher temperatures (approximately 380 K) by increasing the values of both band gap and hole concentration. The maximum ZT of 1.3 was attained at 380 K with the composition x ¼ 1.7 and TE parameters measured along the same direction. More importantly, the average ZT av of approximately 1.18 was achieved in the range of 300-480 K, indicating potential for the application of this composition in low-temperature power generation.

EXPERIMENTAL PROCEDURES Melting
Commercial high-purity elemental chunks of 99.999% Bi, 99.999% Sb and 99.999% Te were used as raw materials. Appropriate quantities of each were weighed according to the nominal composition of Bi 2 Àx Sb x Te 3 (x ¼ 1.4, 1.5, 1.6, 1.7 and 1.8) and sealed in a quartz tube at 10 À3 Pa. The element mixture was melted at 1023 K for 10 h and cooled in the furnace.

Consolidation of powders and HD
Ingots were ball milled (MM 200, Retsch Gmbh, Haan, Germany) for 20 min at 15 Hz to yield fine powders. These powders were hot pressed into cylindrical shapes in a f10 mm graphite die at 673 K for 30 min under 80 MPa. The initial hot-pressed bulk samples with different Sb content (x) were named HP-Sbx. Subsequently, HD was performed by repressing the HP samples in a larger graphite die with an inner diameter of 16 mm at 823 K for 30 min at the same pressure. This HD approach eventually yielded 16 mm disk-shaped samples, labeled 'HD-Sbx' .

Materials characterization
The phase identification and grain orientation of all the samples were investigated using X-ray diffraction with a Rigaku D/MAX 2500/PC diffractometer (Rigaku Corp, Tokyo, Japan). In-plane thermal diffusivity (D) was measured on a Netzsch LFA 457 laser flash apparatus (Netzsch Gmbh, Selb, Germany) using the method introduced by Xie et al. 38 Figure 1a, Supplementary Information). Specific heat (C P ) was measured on the Netzsch DSC 404C (Netzsch Gmbh) (Supplementary Figure 1b) and density (r D ) was estimated by an ordinary dimension and weight measurement procedure. Thermal conductivity in the in-plane direction was then calculated using the relation k ¼ Dr D C P . In-plane electrical conductivity (s) and the Seebeek coefficient (a) were simultaneously measured on a commercial Linseis LSR-3 system (Linseis Gmbh, Selb, Germany). The Hall coefficient (R H ) was determined at 300 K on a Quantum Design PPMS-9T instrument using a four-probe configuration (Quantum Design Inc, San Diego, CA, USA). The carrier concentration (n H ) and in-plane Hall mobility (m H ) were calculated according to n H ¼ À1/eR H (assuming Hall factor of 1.0) and m H ¼ sR H . Figure 1a shows the in-plane X-ray diffraction patterns of the as-pressed Bi 0.4 Sb 1.6 Te 3 bulk samples before and after HD. All the reflection peaks can be indexed to a rhombohedral phase (JCPDS #49-1713). The (00l) diffraction intensities of the HD sample are much higher than observed the HP samples, revealing the formation of preferred orientation during the HD process. To investigate the texture evolution, the degree of preferred orientation (F) of the (00l)planes was calculated by the Lotgering method. 39 As shown in Figure 1b, the F values increase from 0.14-0.17 in the HP series to 0.20-0.28 in the HD samples. Obviously, HD causes enhanced texture in the bulk samples because of the occurrence of plastic deformation. In our previous study, a similar process yielded reduced textures, which was ascribed to deformation-induced dynamic recrystallization. 7 The possible cause of the difference may be the different degree of deformation applied to the samples.

Texture evolution
It is well known that both plastic deformation and dynamic recrystallization take place in HD process and have opposite impacts on the degree of texture. The deformation strain (e, defined by e ¼ (H 0 ÀH)/H 0, where H 0 and H are the initial thicknesses of the HP samples and the final thicknesses of the HD samples, represents the degree of plastic deformation. A value of e ¼ 0.37 was obtained in the earlier report, 7 whereas a value of e ¼ 0.61 was observed in this work. The larger deformation strain induced in the present work led to the formation of texture as a result of the comparatively stronger role played by plastic deformation than dynamic recrystallization. The F values reported here are higher than those observed in materials prepared using the MA-spark plasma sintering (MA: mechanical alloying) 37 or MA-HP 40 techniques. However, processes such as zone melting accompanied by spark plasma sintering 41 or shear extrusion 42 provide stronger textures, with F values of 0.61 (zone melting) and 0.63 (shear extrusion). Scanning Electron Microscopy (SEM) fractographs of the bulk samples give results consistent with X-ray diffraction measurements (Supplementary Figure 2). It should be mentioned that even small degrees of texture will result in large differences between in-plane and out-of-plane thermal conductivities.

Effects of Sb content and HD on TE properties
In p-type Bi 2 Àx Sb x Te 3 alloys, the hole concentration is primarily created by antisite defects in the form of Te sites occupied by Bi or Sb atoms. [25][26][27] The incorporation of Sb atoms into the Bi 2 Te 3 lattice decreases the difficulty associated with the formation of antisite defects and thereby increases the hole concentration. 27,43 Thus, as illustrated in Figure 2a, the hole concentration (n H ) monotonically rises with increasing x. The n H values of Bi 2 Te 3 -based alloys are sensitive to both chemical composition and processing-induced lattice defects. 22,23,29 Tellurium vacancies can be generated by a process of deformationinduced non-basal slip that provides the lattice with excess negative carriers, referred to as the donor-like effect. 28,44 The donor-like effect can partially compensate for the hole concentration in Sb-rich p-type solid solutions. 28 This phenomenon explains the slight reduction in n H observed in the HD samples compared with their HP counterparts (Figure 2a). Single crystals prepared by the Bridgeman method without mechanical deformation show the highest n H . 27,43 The influence of Sb content on the carrier mobility is illustrated in Figure 2b. The markedly increased n H achieved with increasing x leads to strong electron-electron scattering and decreased m H . However, it is interesting that both n H and m H increase with increasing x. The increase of m H with increasing x (a reduced Bi/Sb ratio) is mainly attributed to weakened alloy scattering, which is consistent with previous studies. 24,45 In addition, the HD samples show much higher m H values than the HP samples by virtue of their lower n H and slightly enhanced textures. 22,23,46,47 Nevertheless, the measured m H values of the HD samples are still smaller than the m H for the single crystals 43 and MA-HP 48 samples of the same composition. Further improvements in m H may be possible through the formation of a highly preferred orientation via repetitive HD or optimization of the process parameters.
The electrical transport properties of the Bi 2 Àx Sb x Te 3 bulk samples have been measured and plotted in Figure 3, wherein a transition from semiconducting (x ¼ 1.4) to metal-like (x41.5) conduction behavior is observed. The remarkable enhancement in s with increasing x can be attributed to increases in both n H and m H because of the incremental concentration of antisite defects 27,43 and reduced alloy scattering, 24 respectively. Post-HD changes in s were different for materials with xo1.6 and x41.6 because of the different magnitudes of the n H reduction and the m H increase. In particular, s is improved, to some extent, because of texture enhancement in materials with Sb content x41. 6. The electrical conductivities of the Bi 2 Àx Sb x Te 3 (x ¼ 1.5-1.8) bulk samples roughly follow a T À1.5 dependence near room temperature, indicative of acoustic phonon  Figure 1 (a) Typical X-ray diffraction patterns of the Bi 0.4 Sb 1.6 Te 3 bulk samples before and after hot deformation, measured on the hot pressed surfaces, and (b) F values of the Bi 2 Àx Sb x Te 3 bulk samples before and after hot deformation, as well as data from Chen, 37 Fan, 40 Jiang 41 and Kim. 42 Shifting up of p-type bismuth telluride-based TE materials for power generation L-P Hu et al scattering. The dominant carrier-scattering mechanism does not change after HD processing. Figure 3b displays the temperature dependence of a for bulk Bi 2 Àx Sb x Te 3 samples. Table 1 lists some room temperature physical parameters of the Bi 2 Àx Sb x Te 3 alloys. The reduced Fermi levels (x F ) and carrier effective mass (m*) were roughly estimated from the measured a and n H based on a single parabolic band model with acoustic phonon scattering. 49 Although the calculation using the single parabolic band model may introduce error because of the non-parabolic nature of the valence band and complex scattering processes, the trends in the values of x F and m* can provide information about the band structure upon the introduction of Sb. The increase in n H with increasing x is accompanied by a gradual increase in x F , indicating the progressive movement of the Fermi level deeper into the valence band. However, the Fermi level slightly shifts toward the valence band edge after HD, as evidenced by the decreased x F of the HD samples in comparison to their HP counterparts. m* remains nearly unchanged with increasing x. The slight decrease in m* at x ¼ 1.8 can be attributed to variations in the valence band constituents, consistent with Stordeur's results. 50 Deformation does not cause changes in m*.
The room temperature value of a is reduced from 255 mVK À1 (HD-Sb1.5) to 139 mVK À1 (HD-Sb1.8) with increasing x, because of the increase in n H . The noticeable improvement in a achieved after HD is a result of the reduction in n H . Therefore, HD processing enables the improvement of a through n H reduction, whereas slightly increasing s by enhancing the textures of the Bi 0.3 Sb 1.7 Te 3 samples. The band gaps of the hot-deformed Bi 2 Àx Sb x Te 3 (x ¼ 1.4-1.7) bulk samples were roughly estimated using the relationship between the peak a and the corresponding temperature, that is, E g ¼ 2ea max T max . Slight band gap enhancement was observed from x ¼ 1.4 (E g ¼ 0.11 eV) to x ¼ 1.7 (E g ¼ 0.18 eV). The behaviors of T max and n H as functions of x are also shown in Figure 3c. It is clear that the temperature at which a max is observed gradually rises as the Sb content increases from 1.4 to 1.8 because of the suppression of the detrimental effects of minority carriers on a.
A plot of the calculated power factors PF ¼ sa 2 is shown in Figure 4. The value of PF increases considerably with increasing x, because of a substantial improvement in s. All the HD samples exhibit higher PF values than their HP counterparts, except when x ¼ 1.4. A maximum value of 4.5 Â 10 À3 Wm À1 K À2 was achieved at room temperature for the HD-Sb1.7 sample, approximately 20% higher than without HD.
In-plane thermal conductivity was measured using the method introduced by Xie et al. 38 The variation of in-plane k with temperature is plotted in Figure 5a. The rise of k at room temperature with increasing x is mainly ascribed to an increase in the electronic contribution k el , which is estimated by the Wiedeman-Franz relationship (k el ¼ L 0 sT) using the Lorentz constant, L 0, of 2.0 Â 10 À8 V 2 K À2 . The value of k el calculated at 298 K for the HD-Sb1.4  sample is 0.08 Wm À1 K À1 . This value increases to 1.22 Wm À1 K À1 for the HD-Sb1.8 sample. The value of k ph , approximately equal to k Àk el before the intrinsic excitation, is plotted as a function of temperature in Figure 5b. The HD samples have enhanced (00l) texture; therefore, higher in-plane k values are generally expected. However, the HD samples have lower values of k when compared with their HP counterparts, likely as a result of enhanced phonon scattering by local nanostructures and high-density lattice defects created by HD processing, as demonstrated earlier. 7,[21][22][23] In particular, the extremely low k ph obtained for the HD-Sb1.8 sample (0.28 Wm À1 K À1 ) at 355 K is only approximately 45% of that for its HP counterpart.

Suppression of ambipolar thermal conductivity
Within the intrinsic conduction region, there is an additional component to the thermal conductivity arising from the diffusion of electron-hole pairs, called the ambipolar component k amb . Sizeable values of k amb have serious adverse effects on the k, which are particularly notable in the samples with low Sb content. We evaluated k amb for the materials presented in this work to compare the ambipolar contribution to k for alloys of different Sb content. In cases where k amb has no role, the relationship between the k ph and the reciprocal of temperature (1/T) is theoretically linear (k ph p1/T).
Assuming a negligible effect of k amb on k before intrinsic excitation, the difference between the calculated k ph and experimental k Àk el is approximately equal to k amb (Figure 6a). The value of k amb increases with temperature, but the magnitude of the increase reduces with increasing x. This may be ascribed to the suppression of intrinsic   Shifting up of p-type bismuth telluride-based TE materials for power generation L-P Hu et al conduction by increases in both the hole concentration and band gap (Figure 6b). For example, the k amb value calculated at 400 K for the HD-Sb1.5 sample is 0.36 Wm À1 K À1 , but it is only 0.07 Wm À1 K À1 for the HD-Sb1.8 sample. This approximate 80% reduction indicates that the detrimental effects of intrinsic conduction are substantially suppressed by increased Sb content.
The materials parameter b and dimensionless figure of merit ZT The materials parameter (b, Figure 7) is proportional to (m/k ph ) (m*/m 0 ) 3/2 and determines the TE efficiency. 51 Enhancing TE performance requires an increased ratio of carrier mobility to lattice thermal conductivity, along with a high effective mass. Obviously, the value of (m/k ph )(m*/m 0 ) 3/2 at 300 K rises with increasing Sb content in our bulk polycrystalline samples. HD remarkably improves b, showing its promise as a preparation technique for high performance TE materials. The dimensionless figure of merit, with both electrical and thermal properties measured along the in-plane direction as a function ZT values at 300 and 380 K as a function of Sb content x, as well as data from Ivanova, 53 Yim, 24 Fan. 48   Shifting up of p-type bismuth telluride-based TE materials for power generation L-P Hu et al of temperature, is calculated and displayed in Figure 8a. Considerable enhancement in ZT was obtained via raising x, especially at elevated temperatures. Deformation process further optimizes the ZT. The Sb-rich HD-Sb1.7 sample shows the highest ZT (approximately 1.3) at 380 K, almost 45% better than its HP counterpart, indicating its suitability for low-temperature power generation. Traditional Bi 0.5 Sb 1.5 Te 3 alloys grown by zone melting display a maximum ZT of approximately 1 and optimal hole concentrations near room temperature. 24,52 These values are greatly reduced by the donor-like effects induced by ball milling and HD, so the optimum composition of the polycrystalline bulk samples for power generation at higher temperatures will be richer in Sb (Figure 8b).
Owing to the gradual increases in the band gap and hole concentration with increasing Sb content, the temperature of the maximum ZT shifts gradually to higher temperatures as a result of the suppression of intrinsic conduction. More importantly, there is a significant improvement in the average ZT av value throughout the temperature range studied. The average ZT av value for the HD-Sb1.7 sample is 1.18 (Figure 8c). To corroborate the present results, an HD Bi 0.3 Sb 1.7 Te 3 sample was re-prepared twice using the same procedure. Highly repeatable, excellent ZT values (1.26-1.32) were obtained (Figure 8d). The maximum ZT in this work is higher than values for samples prepared by a microwave-stimulated wet-chemical method (ZT ¼ 1.1) 8 or the Bridgman technique (ZT ¼ 1.0), 24 and is comparable to a value achieved with a bottom-up ball milling approach (ZT ¼ 1.4), 6 as shown in Figure 8e. Li et al. 53 also reported a high ZT (approximately 1.33) at 373 K for ball-milled Bi 0.3 Sb 1.7 Te 3 alloys with SiC nanoparticles but not for Bi 0.5 Sb 1.5 Te 3 . It is also worth noting that the use of both the in-plane s and out-ofplane k to determine the ZT would lead to an overvalued figure of 1.9 (the shadow part in Figure 8e), despite its weak texture. However, property anisotropies were neither mentioned nor discussed in detail for the high ZT values reported in previous studies. 4,6,8 In summary, we successfully shifted the maximum ZT values of p-type bismuth telluride-based alloys to relatively high temperatures (approximately 380 K) by Sb-alloying. The detrimental effects of minority carriers on the Seebeck coefficient and thermal conductivity were suppressed as a result of increases in both the hole concentration and band gap. High-performance bismuth telluride-based alloys were fabricated utilizing the HD method. The formation of preferred orientation during the HD process is beneficial to the improvement of electrical conductivity. A donor-like effect partially compensated for the hole concentration, increasing the Seebeck coefficient. Aside from this, the lattice thermal conductivity was considerably reduced by the presence of recrystallization-induced in situ nanostructures and highdensity lattice defects. As a consequence of these factors, the hotdeformed Bi 0.3 Sb 1.7 Te 3 alloy, rather than the typical Bi 0.5 Sb 1.5 Te 3 composition, showed a maximum ZT of 1.3 at 380 K and the largest average ZT av of 1.18 in the range of 300-480 K. These data indicate significant promise for these materials in low-temperature power generation. As these results proved extremely reproducible, this is a promising approach for the mass production of high-performance TE materials.