Enhanced thermoelectric performance in n-type Bi2Te3-based alloys

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Enhanced thermoelectric performance in n-type Bi2Te3based alloys via suppressing intrinsic excitation Feng Hao, Tong Xing, Pengfei Qiu, Ping Hu, Tianran Wei, Dudi Ren, Xun Shi, and Lidong Chen ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b06533 • Publication Date (Web): 06 Jun 2018 Downloaded from http://pubs.acs.org on June 6, 2018

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ACS Applied Materials & Interfaces

Enhanced Thermoelectric Performance in n-Type Bi2Te3-Based Alloys via Suppressing Intrinsic Excitation Feng Hao,# ‡ Tong Xing,#†, § Pengfei Qiu,*† Ping Hu,†, § Tianran Wei,† Dudi Ren,† Xun Shi,*† Lidong Chen†

†

State Key Laboratory of High Performance Ceramics and Superfine Microstructure,

Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai 200050, China ‡

Semiconductor Manufacturing International Corporation, 18 Zhangjiang Road,

Shanghai 201203, China §

University of Chinese Academy of Sciences, 19 Yuquan Road, Beijing 100049, China

#The two authors contribute equally to this work. Abstract Currently, the application of thermoelectric power generators based on Bi2Te3-based alloys for the recovery of low-quality waste heat is still limited due to the material’s aggravated intrinsic excitation at elevated temperatures. In this study, excessive Te and dopant I are introduced to the n-type Bi2Te2.4Se0.6 material with the purpose to suppress its intrinsic excitation and improve the thermoelectric performance at elevated temperatures. These Te and I atoms act as electron donors to effectively reduce the density of minority carriers (holes) and weaken their negative contribution to the Seebeck coefficient. Likewise, the initial band structure and the carrier scattering mechanism are scarcely altered. Similar with the p-type Bi2Te3-based alloys, we found the “conductivity-limiting” mechanism is also well obeyed in the present n-type Bi2Te2.4Se0.6-based materials. The reduced minority carrier partial electrical conductivity in these Te-excessive and I-doped Bi2Te2.4Se0.6 samples significantly decreases the bipolar thermal conductivity, leading to the lowered total thermal conductivity at elevated temperatures. Finally, the peak zT is successfully shifted up to the higher temperatures for these Te-excessive and I-doped Bi2Te2.4Se0.6 samples. A maximum zT of 1.0 at 400 K and average zT of 0.8 at ACS Paragon Plus Environment

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300-600 K have been realized in Te-excessive Bi2Te2.41Se0.6. Keywords: thermoelectric; bismuth telluride; intrinsic excitation; minority carrier; power generation 1. Introduction Thermoelectric (TE) technology can directly convert waste heat into useful electrical power.1 Due to the advantages of high reliability, vibration-free operation, and environmental friendliness, the TE technology has been extensively used in the fields such as space power generation and portable beverage coolers since the mid-20th century. Recently, driven by the global energy crisis and environmental concern, the application of the TE technology is being extended to the field of power generation used in industry such as the recovery of waste heat exhausted from the vehicles or industries. The energy conversion efficiency of a TE material is governed by the dimensionless figure of merit zT (zT = S2σ/κ), where S is the Seebeck coefficient, σ is the electrical conductivity, T is the absolute temperature, and κ is the thermal conductivity.2 For a narrow band gap semiconductor, κ is usually comprised of three parts, termed as the lattice part (κL), carrier part (κe), and bipolar part (κb).3 An ideal TE material should have excellent electrical transport properties (high S and large σ) as well as poor thermal transport properties (low κ). Despite the continuously updated zT records in various novel high performance TE materials in recent years, classic bismuth telluride (Bi2Te3) based alloys are still one of the most excellent TE materials around room temperature.4 Bi2Te3 is a typical narrow-gap semiconductor which has been studied extensively since the 1960s. It possesses a rhombohedral structure (R3m) consisting of three 2D quintuple layers (Te1–Bi–Te2–Bi–Te1).5 These 2D layers are weakly bound to each other by van der Waals force, forming distinct electrical and thermal transport properties along and perpendicular the basal plane. Commercially, the Bi2Te3-based alloys prepared by the zone-melting (ZM) method, with the peak zTs around 300 K, have already realized the large-scale utilization for TE cooling applications. Recently, great efforts in TE

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community have been paid on increasing the zTs of polycrystalline Bi2Te3-based alloys by using the nanostructure engineering.6-14 Via introducing numerous grain boundaries or nanoprecipitates to strengthen the scattering to the heat-carrying phonons, the lattice thermal conductivities of polycrystalline Bi2Te3-based alloys can be significantly suppressed to yield higher zTs than those prepared by the ZM method.15-18 However, above room temperature, the zTs for most of the Bi2Te3-based alloys quickly degrade.7,15 Thus, the application of Bi2Te3-based alloys in TE power generation to harvest low temperature waste heat, especially for the ones during 100-300 ˚C, is still limited so far. The quickly degraded zTs for Bi2Te3-based alloys at elevated temperatures are mainly attributed to the intrinsically excited electron-hole pairs across the band gap.19 For Bi2Te3, the band gap (Eg) is ~0.2 eV, only several kBT above room temperature. Thus, the electrons in the valence band easily acquire enough energy to be thermally excited into the conduction band, leaving same amount of holes in the valence band. These electron-hole pairs not only contribute additional bipolar thermal conductivity κb to the total heat conduction, but also reduce the total Seebeck coefficient, yielding low zTs at elevated temperatures.20-22 Thus, substantially suppressing the intrinsic excitation is critical to the real applications of Bi2Te3-based alloys in TE power generation. Towards this goal, enlarging the Eg to increase the energy required for the electrons to be excited from the valence band to the conduction band is one effective method. For instance, via enlarging the Eg from 0.11 eV to 0.18 eV by Sb-alloying, the peak zT of p-type Bi2Te3-based alloys can be shifted up from room temperature to higher temperature.23-25 Similar strategy has been also performed in n-type Bi2Te3-based alloys by alloying Se at Te-sites.26,27 Beyond enlarging the Eg, the intrinsic excitation can be also suppressed by decreasing the minority carriers partial electrical conductivity (σm) at elevated temperatures.28 In our previous work, we successfully reduced the σm by introducing a tiny amount of electron acceptors (Cd, Cu, and Ag) in p-type Bi2Te3-based alloys.29-31 Correspondingly, the κb at elevated temperatures is significantly lowered, leading to a peak zT of 1.4 at 450 K. However, so far, the related investigation in n-type Bi2Te3-based alloys is still absent. ACS Paragon Plus Environment


In this study, a tiny amount of excessive Te and I dopant is used as electron donors to improve the TE performance of n-type Bi2Te2.4Se0.6 alloys at elevated temperatures. These electron donors significantly reduce the σm, and thereby lessen the negative effect of intrinsic excitation on the electrical and thermal transports. Likewise, it is found that the bipolar thermal transports in these n-type Bi2Te2.4Se0.6 alloys also obey the “conductivity-limiting” mechanism. Due to the substantially suppressed intrinsic excitation, the peak zTs of these n-type Bi2Te2.4Se0.6 alloys are successfully shifted up from room temperature to the higher temperatures. A maximum zT of 1.0 at 400 K and average zT of 0.8 at 300-600 K have been realized in Te-excessive Bi2Te2.41Se0.6. 2. Experimental section High purity elements Bi (99.999%), Te (99.999%), and Se (99.999%), and dopant TeI4 (99.999%) are weighed out according to the stoichiometric ratio of Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02). Then, they were sealed into quartz tubes in vacuum and melted at 1173 K for 12 hours. Subsequently, these quartz tubes were quenched into cold water, following with the annealing process at 673 K for 48 hours. The obtained ingots were ground into fine powders at agate mortar and sintered into dense bulk sample by spark plasma sintering (SPS) technique. During sintering process, the samples were kept at 693 K under 50 MPa for 10 min. Then, the sintered bulk samples were sealed into quartz tubes in vacuum and annealed at 723 K for extra 7 days to obtain the final products. The densities of all products are above 95% of the theoretical value. The heights of sintered cylinders are about 12 mm, which facilitate both the electrical and thermal transport properties characterizations in the directions along and perpendicular to the sintering press. X-ray diffraction (XRD, D8 ADVANCE, Bruker Co. Ltd) was used to analyze the material’s phase purity. The morphology of fractured surface and element distribution were characterized by scanning electronic microscopy (SEM, Supra 55, ZEISS) equipped with an energy dispersive spectrometer (EDS, Oxford Horiba 250). The pole ACS Paragon Plus Environment

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figures were plotted by the analysis software (HKL-Channel 5, Oxford) based on the data of Euler Angle of crystal, which was measured via back scattering electron diffraction (EBSD, Oxford) detector equipped at Magellan 400 field emission scanning electron microscopy (FESEM). which was collected via back scattering electron diffraction (EBSD) detector equipped at Magellan 400 field emission scanning electron microscopy (FESEM). The thermal conductivity was calculated by multiplying the thermal diffusivity (D), heat capacity (Cp) and density (ρ). The D was measured by LFA457 manufactured by Netzsch Co. Ltd according to the laser flash method. The Cp was estimated by using the Neumann-Kopp law.32 The ρ was measured by using the Archimedes method. The σ and S above room temperature were measured simultaneously in ZEM-3 apparatus (ULVAC Co. Ltd). The TE transport properties, including the Hall coefficient (RH), σ, and κ below room temperature, were characterized by the physical property measurement system (PPMS, Quantum Design). The low-temperature RH was measured by sweeping the magnet field to 3 T along both positive and negative directions. The carrier concentration n can be derived with relation n=1/eRH, and then the mobility µH was calculated according to the relation µH = σRH. 3. Results and discussion Figure 1a shows the X-ray diffraction (XRD) patterns for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) powder samples before spark plasma sintering. All diffraction peaks can be indexed to the hexagonal Bi2Te3, but the peak positions slightly shift to the high-angle direction due to the smaller atomic radius of Se (1.16 Å) as compared with Te (1.36 Å). No secondary phases are observed for all samples, indicating that all the excessive Te and the doped I might enter into the lattice of Bi2Te2.4Se0.6. Moreover, all diffraction peaks are very sharp, suggesting that the prepared samples possess good crystallinity. Due to the inherent layered crystal structure, the grains of Bi2Te3-based alloys prefer to be oriented along the crystallographic (00l) planes during the sintering process. This can be confirmed by the distinct XRD patterns performed on the ACS Paragon Plus Environment


sintered Bi2Te2.4Se0.6 bulk sample in the directions along and perpendicular to the sintering press. As shown in Figure 1b, the intensity of the (006) diffraction peak at 2θ = 17º in the direction perpendicular the sintering press is much stronger than that along the sintering press. The orientation factor (F) is usually used to characterize the degree along the (00l) planes in Bi2Te3-based bulk alloys. The F value is 0.34 in the direction perpendicular the sintering press, about 3 times of that along the sintering press. Such highly orientation can be also reflected by the SEM characterization. As shown in Figure 2a, the presence of large grains with layered morphology is observed in the sintered Bi2Te2.4Se0.6 bulk sample. The pole figure obtained from the EBSD characterization performed in a large area (259 × 177 µm2) further proves that the grain orientation in the sintered Bi2Te2.4Se0.6 bulk sample is highly centralized along the (00l) planes (see Figure 2b). With the purpose to maximize the electrical transport performance, all the following TE properties measurements were performed in the direction perpendicular to the sintering press.

Figure 1 (a) X-ray diffraction patterns for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) powder samples. (b) X-ray diffraction patterns for the sintered Bi2Te2.4Se0.6 bulk sample in the directions along and perpendicular to the sintering press.


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Figure 2 (a) Fractured morphologies of the sintered Bi2Te2.4Se0.6 bulk sample. (b) Pole figure about (00l) planes for the sintered Bi2Te2.4Se0.6 bulk sample. The number 1-5 represents the strength of orientation.

Figure 3 shows the EDS elemental mapping for the sintered Bi2Te2.4Se0.6 bulk sample. All the three elements, Bi, Se, and Te, are homogeneously distributed. No obvious secondary phases and elemental segregation are observed. In fact, experimentally, it is more difficult to realize the homogeneous elemental distribution in n-type Bi2(Se,Te)3 alloys than that in p-type (Bi,Sb)2Te3 alloys. Because the melting point for Bi2Se3 (701 ºC) is higher than that for Bi2Te3 (585 ºC),33 the elemental segregation is easy to appear in the n-type Bi2(Se,Te)3 alloys when they are cooled from high temperature to room temperature. Such elemental segregation would severely deteriorate the carrier mobility and degrade the zT. In the present study, an extra annealing procedure is used after the SPS sintering process. This procedure provides the further element diffusion among the grains with different chemical compositions. Thus, high quality n-type Bi2(Se,Te)3 samples with homogeneously distributed Bi, Se, and Te elements are obtained in the present study (see Figure 3).



Figure 3 Backscattered electron (BSE) imaging map and elemental distribution in the sintered Bi2Te2.4Se0.6 bulk sample.

Figure 4a-b shows the electrical conductivity (σ) and Seebeck coefficient (S) for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) samples. All samples possess negative S throughout the entire measured temperature range, indicating that electrons are the dominant carriers. For the pristine Bi2Te2.4Se0.6, such n-type electrical transport behavior is attributed to the formation of Te or Se vacancies during the fabrication process.6 The σ for the pristine Bi2Te2.4Se0.6 is only 2.5×104 Sm-1 at 300 K. With increasing the excessive Te content or the doped I content, the σ gradually increases in the whole measured temperature range. When x = 0.02, the σ is 2.6×105 Sm-1, about one order of magnitude higher than that for the pristine Bi2Te2.4Se0.6. Correspondingly, below 400 K, the S decreases with increasing the excessive Te content or the doped I content. When x = 0.02, the S is 76 µVK-1 at 300 K, about one fourth of that for the pristine Bi2Te2.4Se0.6. However, above 400 K, the S for the pristine Bi2Te2.4Se0.6 quickly degrades with increasing temperature. At 600 K, its S is even lower than all of the other samples. In addition, it seems that the modification of S and σ by doped I is stronger than that by excessive Te. As shown in Figure 4a-b, when the excessive Te content δ equals to the doped I content x, the


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I-doped samples always possess the larger σ but smaller S than the Te-excessive samples. Figure 4c illustrates the temperature dependence of power factor (PF = S2σ) for all samples. Due to the enhanced σ, the PFs are significantly improved in the entire measurement temperature range. At 300 K, the PF for Bi2Te2.41Se0.6 is 31.3 µWcm-1K-2, which is about 50% higher than the pristine Bi2Te2.4Se0.6. Even at elevated temperatures, such as 600 K, the high PF for Bi2Te2.41Se0.6 is still maintained with a value of 13.7 µWcm-1K-2, which is about two times of that for the pristine Bi2Te2.4Se0.6. The thermal conductivity (κ) as a function of the temperature for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) samples is shown in Figure 4d. Introducing the excessive Te or doping with I greatly raises the κ at 300 K. For Bi2Te2.42Se0.6, the κ is 2.2 Wm-1K-1 at 300 K, about 120% enhancement as compared with that for the pristine Bi2Te2.4Se0.6. In Bi2Te3-based alloys, the κ is more complex than that in the wide-gap semiconductors, such as SiGe34, BiCuSeO35, and CuInTe2/CuGaTe236. Besides the lattice thermal conductivity κL and the electronic thermal conductivity κe, the bipolar thermal conductivity κb must be also considered because it weighs a large proportion in the total κ, especially at elevated temperatures.22 Here, the κ increments around room temperature are mainly contributed by the enhanced electronic thermal conductivity κe. Generally, κe can be calculated using the Wiedeman-Franz law (κe = LσT, where L is the Lorenz number, L = 1.6×10-8 V-2K-2 in the present study).37 As shown in Figure 4e, introducing the excessive Te or doped I greatly raises the κe values around the room temperature due to the significantly enhanced σ (see Figure 4a). Above 500 K, although the κe is also increased, the total κ of the samples including excessive Te or doped I are lower than that of the pristine Bi2Te2.4Se0.6. This is attributed to the reduced bipolar thermal conductivity κb, which will be discussed later. For Bi2Te2.42Se0.6, κ is 1.6 Wm-1K-1 at 600 K, about 30% lower than pristine Bi2Te2.4Se0.6. Such lowered κ and the well maintained high PF simultaneously lead to the obviously improved zTs at elevated temperatures. As shown in Figure 4f, the peak zT of 0.6 appears at 320 K for the pristine Bi2Te2.4Se0.6. With increasing the excessive Te content or the doped I content, ACS Paragon Plus Environment


the peak zT gradually shifts to the high temperatures. For the Te-excessive Bi2Te2.41Se0.6, the peak zT of 1.0 appears at 420 K. Moreover, the average zT for Bi2Te2.41Se0.6 is 0.8, about 116% enhancement as compared with the pristine Bi2Te2.4Se0.6.

Figure 4 Temperature dependences of (a) electrical conductivity σ, (b) Seebeck coefficient S, (c) power factor PF, (d) thermal conductivity κ, (e) electronic thermal conductivity κe, and (f) figure of merit zT for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) samples.

Carrier concentration is an important parameter to analyze the TE properties. Because I has one more electron in the outmost shell than Te, doping I in Bi2Te2.4Se0.6 would increase the electron concentration. Meanwhile, introducing the excessive Te ACS Paragon Plus Environment

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increases the density of Te• antisite defects inside the lattice.38 These Te• antisite defects act as the electron donors, which would also increase the electron concentration. Figure 5a presents the measured Hall carrier concentration (n) for all Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) samples. As expected, n increases in the entire temperature range with increasing the excessive Te content or the doped I content. When x = 0.02, n is 1.2×1020 cm-3 at 300 K, about one order of magnitude higher than the pristine Bi2Te2.4Se0.6. This variation range is much broader than that achieved by just adjusting the Te/Se atomic ratio, such as from 6.0×1019 cm-3 for Bi2Te2Se1 to 1.0×1020 cm-3 for Bi2Te1.1Se1.9.26 Likewise, when δ = x, the Te-excessive samples always possess the lower n than the I-doped samples, which can well explain the smaller σ but larger S in the Te-excessive samples than those in the I-doped samples. The relative lower n in Te-excessive Bi2Te2.4Se0.6 samples should be attributed by the fact that only part of Te atoms enter into the Te-sites forming the Te• antisite defects.

The S and σ variations shown in Fig. 4a-b suggest that the contribution from minority carriers (holes) in the room-temperature electrical transports can be neglected for the current n-type Bi2Te3-based materials. Thus, here we use the single parabolic band (SPB) model to analyze the electrical transports at 300 K for all Bi2Te2.4Se0.6-based samples. In SPB model, the Seebeck coefficient S and carrier concentration n satisfy the following equations: 39,40

=−

(⁄⁄ ( ( − (⁄⁄ ( *

( = !+

"# $"

%&'( (")

⁄ 0∗ 2

, = 4. /

3

4

%⁄ (

(1) (2) (3)

Herein, kB is the Boltzmann constant, e is the elementary charge, h is the Planck constant, ηe is the reduced Fermi level for electrons, and λ is the scattering factor that could be taken as -1/2 for the scattering mechanism including both electron-phonon scattering and alloy scattering. When m* is taken as 1.15me (where me is the inertia mass of free electron), the calculated curve agrees well with the experimental data (see Figure 5b). This suggests that introducing excessive Te or doping with I in the ACS Paragon Plus Environment


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n-type Bi2Te2.4Se0.6 alloys scarcely changes the initial band structure near the Fermi level. This is reasonable considering the tiny amount of Te and I used in the present study. Figure 5c depicts the Hall mobility (µH) for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) samples below 300 K. The µH for the pristine Bi2Te2.4Se0.6 is 192 cm2V-1s-1 at 300 K, much higher than the values for the same composition reported in the previous literatures, such as 141 cm2V-1s-1 by Zheng et al.41 The highly orientation and homogeneous element distribution mentioned above are responsible for the high µH achieved in the present study. Introducing excessive Te or doping with I significantly reduces the µH in the entire measured temperature range. When the I-doped content x = 0.02, the µH is 135 cm2V-1s-1 at 300 K, about 70% of the value for the pristine Bi2Te2.4Se0.6. However, the temperature dependence of the µH is scarcely changed by introducing excessive Te or doping with I. This is reasonable because the Bi2(Te,Se)3 samples usually possess large dielectric permittivities (ε = 80 for Bi2Te3 and ε = 100 for Bi2Se3) and small effective masses (~1.2me), which can effectively suppress the long-range Coulomb potentials and lessen the effect of ionized impurities on the electrical transports.42-44 The weak influence of the excessive Te or doped I on the carrier scattering mechanism can be also reflected by the µH ~ n relationship. Under the assumption that the carrier transport at room temperature is only dominated by the acoustic phonon scattering, the theoretical µH ~ n relationship can be predicted by the following equations45

5=

√ℏ8 9:

(0 ∗ @⁄ ( 2⁄ ; A ?

B⁄ ( C (

5D = ED 5

(4) (5)

Here, ℏ is the reduced Planck constant, ρ is density, ν is longitudinal sound speed (taken as 2884 m/s) 46, FG∗ is the single valley effective mass and ED is taken as 1 for the semiconductors dominated by the acoustic phonon scattering.47 The deformation potential EAC reflects the intensity of interaction between electrons and phonons.48 As shown in Figure 5d, when EAC = 9.5 eV and m* = 1.15 me, the theoretical µH ~ n relationship can well interpret the experimental data, further proving


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that the excessive Te or doped I in Bi2Te2.4Se0.6 only introduces additional carriers while scarcely changes the initial carrier scattering mechanism.

Figure 5 (a) Temperature dependence of carrier concentration n, (b) carrier concentration n dependence of Seebeck coefficient S, (c) temperature dependence of carrier mobility µH, and (d) carrier concentration n dependence of carrier mobility µH for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) samples.

Based on the measured n, the minority carrier concentration (p) and their partial electrical conductivity (σm) can be calculated. Based on the reduced Fermi level η calculated by Equation 1, the reduce Fermi energy for holes ηh can be obtained by subtracting η from -Eg/kBT (where Eg is about 0.2 eV for the pristine Bi2Te2.4Se0.6). Then, the hole concentration p at 300 K is40 ⁄ 0I ∗ 2

H = 4. /

3

4

%⁄ (3

(6)

where mh* is the effective mass of the holes (mh* = 1.1 me).28 In semiconductors, the product of electron concentration (n) and hole concentration (p) at temperature T can be expressed as49 ACS Paragon Plus Environment


,(J ∙ H(J = LJ exp(−OP /RS T

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(7)

where A is a temperature-independent constant under the assumption that electronic structure is temperature-independent. Then, by using the estimated p and experimental n data at 300 K, the constant A in Equation 7 can be calculated. At temperatures above 300 K, the intrinsic excitation occurs. Electrons are excited from valence band to conduction band, leaving the same amount of holes in valence band. Thus Equation 7 is rewritten as (,++T + ∆ ∙ (H++T + ∆ = LJ exp(−OP /RS T

(8)

where ∆ represents the excited concentration of electron-hole pairs at temperature T. In this way, the actual n and p values at specified temperature above 300 K can be obtained. As shown in Figure 6a, p for the pristine Bi2Te2.4Se0.6 is about 3.6×1016 cm-3 at 300 K, which is about 3 orders of magnitude lower than the electron concentration n. However, it quickly increases when increasing temperature due to the aggravated intrinsic excitation, contributing more to the Seebeck coefficient at elevated temperatures. In comparison, n is weakly dependent on the temperature during 300-600 K, which is presented in Figure 6b. Because the Seebeck coefficient is negative for electrons and positive for holes, the increased p leads to the quickly degraded S at elevated temperatures (see Figure 4b). However, when introducing a tiny amount of excessive Te or doped I, the p values are hugely decreased in the whole temperature range. The p for Bi2Te2.42Se0.6 at 300 K is reduced to only 7% of that of the pristine Bi2Te2.4Se0.6. This is consistent with the weak degradation of the Seebeck coefficient in those Te-excessive and I-doped materials at elevated temperatures (see Figure 4b).


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Figure 6 Temperature dependences of the calculated (a) hole concentration p and (b) electron concentration n for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) samples.

Bipolar thermal conductivity κb is also very important for understanding the intrinsic excitation behavior. At the temperatures higher than the Debye temperature (162 K for Bi2Te3),20,50 the phonon-phonon Umklapp scattering dominates the heat transport, i.e. the temperature dependence of κL satisfies the relation of κL ∝ T -1.50,51 At the temperatures when the contribution from minority carriers can be neglected, the (κ - κe) values should be approximately equal to κL. Thus, via fitting the (κ - κe) data at proper temperature range according to expression κL = aT-1 + b (where a and b are fitting parameters), the κL at higher temperature for the present n-type Bi2Te3-based alloys can be obtained. For the pristine Bi2Te2.4Se0.6, the (κ - κe) data at 200-250 K are used. For other samples, the (κ - κe) data at 300-330 K are used. Taking the pristine Bi2Te2.4Se0.6 as an example, Figure 7a shows that this relation can well fit the (κ - κe) data. Similar good fitting has been also observed for other samples, such as Bi2Te2.39I0.01Se0.6 (see Figure 7b). The fitted parameters for the samples prepared in the present study are listed in Table 1 and the estimated κLs are plotted in Figure 7c. The estimated κL for all samples are quite close at elevated temperatures. Then, the κb can be calculated according to the expression of κb = κ - κe - κL. The calculated data are presented in Figure 7d. Clearly, the samples with excessive Te or I dopants possess much lower κb. When the content of I dopants is 0.02, the κb at 600 K is only



0.18 Wm-1K-1, about 87% lower than the pristine Bi2Te2.4Se0.6. Such significantly reduced κb is also one main reason for the improved zT at elevated temperature.

Figure 7 Fitted (κ - κe) curves according to the relation of κL = aT-1 + b for (a) Bi2Te2.4Se0.6 and (b) Bi2Te2.39I0.01Se0.6. Temperature dependences of (c) lattice thermal conductivity κL and (d) bipolar thermal conductivity κb for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) samples.

Table 1 Fitting parameters for all samples in Figure 7a using the expression of κL = aT-1 +b for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) samples. κL = aT -1 + b a b

δ=0 234.29 0.08

Bi2Te2.4+δSe0.6 δ = 0.01 δ = 0.02 160.00 212.26 0.24 0.11

Bi2Te2.4-xIxSe0.6 x = 0.01 x = 0.02 274.28 289.12 0.09 0.04

In the extrinsic semiconductors, Wang et al. suggested that the bipolar thermal transport is in general a “conductivity-limiting” mechanism, and it is controlled by the


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minority carrier partial electrical conductivity.28 In p-type Bi2Te3-based alloys, this mechanism has been already proved.29,30 In the present n-type Bi2Te3-based alloys, we found this mechanism also works. Figure 8a shows the minority carrier partial electrical conductivity σm for the present n-type Bi2Te3-based alloys by subtracting the electron partial electrical conductivity (W ) from the total electrical conductivity σ. The σe is evaluated according to the relation W = X,5 , where 5 is the electron mobility. The 5 above 300 K can be obtained by extrapolating the experimental mobility data in low temperature range of 200-300 K based on the relationship of 5 (J ~ J )/. Clearly, the σm decreases with increasing the excessive Te content or the I-doped content. At elevated temperatures, the σm decrement is more obvious. The σm for Bi2Te2.38I0.02Se0.6 at 600 K is reduced to only 25% of that for the pristine Bi2Te2.4Se0.6. In Figure 8b, we plot the relationship between the κb and σm at 400-600 K for the Bi2Te2.4+δSe0.6 (δ = 0, 0.01, 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01, 0.02) samples. At each specified temperature, κb monotonously increases with increasing σm, confirming the validity of the “conductivity-limiting” mechanism in the present n-type Bi2Te3-based alloys.28 This further proves that introducing the electron donors to suppress minority carrier partial electrical conductivity is a very effective approach to suppress the intrinsic excitation in the present n-type Bi2Te3-based alloys.

Figure 8 (a) Temperature dependence of hole partial electrical conductivity σm for excess Te or I doped Bi2Te2.4Se0.6 samples. (b) Bipolar thermal conductivity κb as a function of hole partial electrical conductivity σm at 400-600 K for the Bi2Te2.4+δSe0.6



(δ = 0, 0.01, and 0.02) and Bi2Te2.4-xIxSe0.6 (x = 0.01 and 0.02) samples. 4. Conclusions In summary, we have successfully optimized the thermoelectric performance of n-type Bi2Te2.4Se0.6 alloys at elevated temperatures by introducing excessive Te or doping with I at Te sites. These Te and I atoms act as the electron donors and significantly decrease the hole concentration to suppress the intrinsic excitation, which is well reflected by the reduced bipolar thermal conductivity and the decreased minority carrier partial electrical conductivity. However, the initial band structure and the carrier scattering mechanism are scarcely altered. Finally, it is found that the bipolar thermal transports in these n-type Bi2Te2.4Se0.6 alloys also obey the “conductivity-limiting” mechanism. As a result of the suppressed intrinsic excitation, a maximum zT of 1.0 at 400 K and average zT of 0.8 at 300-600 K have been realized in Te-excessive Bi2Te2.41Se0.6. Author Information Corresponding Authors *E-mail: [email protected]. *E-mail: [email protected]. ORCID Pengfei Qiu: 0000-0001-6011-1210 Xun Shi: 0000-0002-8086-6407 Tong Xing: 0000-0002-4548-7209 Notes The authors declare no competing financial interest. Acknowledgement This work is supported by the National Natural Science Foundation of China (NSFC) under the No. 51625205, the Key Research Program of Chinese Academy of Sciences


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(Grant No. KFZD-SW-421), the International S&T Cooperation Program of China (2015DFA51050),

and the

Program

of Shanghai Subject Chief

Scientist

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Enhanced thermoelectric performance in n-type Bi2Te3-based alloys

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