Tellurium doped n-type Zintl Zr3Ni3Sb4 thermoelectric materials: Balance between carrier-scattering mechanism and bipolar effect

Materials Today Physics xxx (2017) 1e8

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Tellurium doped n-type Zintl Zr3Ni3Sb4 thermoelectric materials: Balance between carrier-scattering mechanism and bipolar effect Zihang Liu a, Jun Mao a, b, Shengyuan Peng c, Binqiang Zhou d, Weihong Gao c, Jiehe Sui e, *, Yanzhong Pei d, **, Zhifeng Ren a, *** a

Department of Physics and TcSUH, University of Houston, Houston, TX, 77204, USA Department of Mechanical Engineering, University of Houston, Houston, TX, 77204, USA School of Materials Science and Engineering, Harbin Institute of Technology, Harbin, 150001, China d School of Materials Science and Engineering, Tongji University, Shanghai, 201804, China e State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin, 150001, China b c

a r t i c l e i n f o

a b s t r a c t

Article history: Received 9 August 2017 Received in revised form 24 August 2017 Accepted 25 August 2017 Available online xxx

Thermoelectric figure of merit ZT has been greatly improved in the past decade via band engineering to enhance power factor or nanostructuring to reduce thermal conductivity, but less attention has been paid to other significant factors, e.g., carrier scattering mechanism, bipolar effect, etc. Here we show that Te doping on the Sb site, as an n-type strong donor, could significantly suppress the high-temperature bipolar effect in the nanostructured Zintl Zr3Ni3Sb4, which can be ascribed to the combination of high majority-carrier concentration and enlarged band gap. A relatively good ZT of ~0.6 at 773 K for Te doping can be achieved and that is almost double of the previous reported ZT by Cu doping. In addition, the role of carrier scattering mechanism on the low-temperature electrical transport properties is also pointed out, where both carrier mobility and power factor of Te doping, due to the detrimental effect of ionized impurity scattering, are lower than that of Cu doping in which the mixed acoustic phonon and ionized impurity scattering dominates. © 2017 Elsevier Ltd. All rights reserved.

Keywords: Thermoelectric materials n-type Zintl Zr3Ni3Sb4 Bipolar effect Carrier scattering mechanism

1. Introduction Thermoelectric power generation as a clean and sustainable energy harvesting technology has received strong interests in the past two decades [1e5]. However, nowadays thermoelectric devices are still restricted in niche markets, primarily due to the relatively low conversion efficiency that is dominated by the materials' thermoelectric figure of merit, ZT ¼ (S2s/ktotal)T, where S, s, ktotal, and T are the Seebeck coefficient, electrical conductivity, total thermal conductivity, and absolute temperature, respectively. Therefore, pursuing higher ZT, especially for higher average ZT, is always the target for thermoelectric community [6e11]. Narrow bandgap semiconductors are generally considered as promising thermoelectric candidates for power generation [12].

* Corresponding author. ** Corresponding author. *** Corresponding author. E-mail addresses: [email protected] (J. Sui), [email protected] (Y. Pei), [email protected] (Z. Ren).

Power factor (PF ¼ S2s) is a critical material-level parameter that determines the final output power density of devices [13,14]. The fundamental parameters (S and s) are intertwined or even contradicted for a given semiconductor, hence it is a longstanding challenge to largely improve the power factor [2,15]. In principle, optimizing carrier concentration n is a universal and effective approach to enhance power factor [16e20]. Since point defects, originated from intrinsic defects and doping effect, are ubiquitous scattering centers for carrier in semiconductor, recently tuning carrier scattering mechanism has become another fundamental degree of freedom to optimize electrical properties [21e23]. For example, Nb-doping in n-type Mg3Sb2-based materials could convert the ionized impurity scattering into mixed scattering, which led to the triple enhancement of room-temperature PF [23]. For thermal part, the total thermal conductivity ktotal includes contributions from lattice klat, electron kele, and bipolar diffusion kbip. Due to the structural dependence of klat, most efforts have been placed on reducing klat by alloying or nanostructuring [24e29]. Moreover, bipolar diffusion at high temperature would make a significant contribution to heat transport that sharply decreases the Seebeck coefficient concurrently [12]. In general, increasing the

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Please cite this article in press as: Z. Liu, et al., Tellurium doped n-type Zintl Zr3Ni3Sb4 thermoelectric materials: Balance between carrierscattering mechanism and bipolar effect, Materials Today Physics (2017), https://doi.org/10.1016/j.mtphys.2017.08.002

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concentration of majority carrier by heavy doping, suppressing the minority carrier mobility by nanostructured barriers or widen the band gap by isoelectronic alloying could suppress the bipolar effect to a certain extent [6,30e35]. Recently, Wang et al. pointed out that bipolar thermal transport is in general a “conductivity-limiting” phenomenon in semiconductors [36]. In a word, minimizing heat transport - “Phonon Glass” while simultaneously preserving good electrical properties - “Electron Crystal” is the main challenge for realizing high ZT [37]. Zintl phases are typical examples for applying this concept to achieve high ZT [38e41]. Zintl phases comprise electropositive cations and the more electronegative non-metallic elements, where cations donate all valence electrons to anions in order to satisfy valence [42]. According to the Zintl-Klemm concept [43], the anion forms a covalently bonded substructure, resulting in the large unit cell and complex crystal structure. Therefore, the most intriguing trait of Zintl phases as thermoelectric materials is the intrinsically low klat and thus proper doping can be adopted to tune the electrical transport properties [40]. Among Zintl phases, Zintl antimonides are the most promising category for intermediate or high temperature power generation, e.g., Yb14MnSb11 [44], AZn2Sb2 (A ¼ Ca, Sr, Yb or Zn) [45,46], Ca3AlSb3 [47], etc. However, the majority of binary and ternary Zintl antimonides are p-type semiconductors that is probably related with the cation vacancy [48], which limit the practical applications owing to lack of suitable ntype counterparts [49]. As a consequence, it is imperative to develop high performance n-type Zintl antimonides. Very recently, a high ZT value of ~1.5 is achieved for the n-type Zintl phase Mg3Sb2-based materials by utilizing defect chemistry and introducing Sb/Bi disorder [50], which opens up a new and effective strategy to design n-type Zintl antimonides. In fact, it is highly beneficial to fabricate a pair of n-type and ptype thermoelectric materials based on one parent compound, because those two candidates would have the similar thermoelectric properties and thermal expansion coefficient [37]. Recently, both p-type and n-type characteristic could be realized by Co doping and Cu doping on the Ni site in the Zr3Ni3Sb4 system [51,52], respectively. Zintl phase Zr3Ni3Sb4 compound exhibits the Y3Au3Sb4-type structure, a filled variant of the more prevalent Th3P4 structure [53], where Ni atoms enter the distorted tetrahedral sites that weakly hybridize with the existing Zr3Sb4 network [54,55]. However, low peak ZT (~0.5 for p-type and ~0.4 for n-type) is the major drawback, where the bipolar effect seriously deteriorates the high-temperature performance. Thus, how to effectively suppress the bipolar effect becomes the primary task for Zr3Ni3Sb4 based materials. Herein, we directly employ mechanical alloying and hot pressing procedure to fabricate the nanostructured Te doped Zr3Ni3Sb4 alloys. A decent peak ZT of 0.6 can be obtained, which can be ascribed to the weakened or vanished bipolar effect at high temperature. In addition, a detailed investigation of the effect of Cu and Te doping on the carrier mobility has been performed to highlight the role of carrier scattering mechanism on the low-temperature electrical transport properties. 2. Experimental section 2.1. Synthesis Zr sponges (99.95%; Atlantic Metals and Alloy), Ni powder (99.9%; Sigma Aldrich), Sb rods (99.999%; Alfa Aesar), and Te chunks (99.99%; Alfa Aesar) were weighed according to the nominal composition Zr3Ni3Sb4-xTex samples (x ¼ 0, 0.025, 0.05, 0.1 and 0.2), loaded into the stainless steel jar in a glove-box under an argon atmosphere and then subjected to ball-milling for 12 h. The ball-

milled powder was loaded into the half-inch die and hot pressed by direct current (dc-HP) press at 1223 K for 2 min under a pressure of ~90 MPa. 2.2. Sample characterization X-ray diffraction (XRD) analysis was performed using a PANalytical multipurpose diffractometer with an X'celerator detector (PANalyticalX'Pert Pro) and lattice parameters was calculated by Rietveld refinement method. The phases were analyzed with JADE 6.0 software. The morphologies and actual composition were analyzed using a Hitachi S4700 scanning electronic microscope (SEM) accompanied with energy-dispersive X-ray spectroscopy (EDS). 2.3. Transport property measurements Bar samples were cut from the disks and used for simultaneous measurement of electrical resistivity (r) and Seebeck coefficient (S) on a commercial system (ULVAC ZEM-3). Thermal conductivity was calculated using k ¼ DCpd, where D, Cp, and d are the thermal diffusivity, specific heat capacity, and density, respectively. The thermal diffusivity coefficient (D) was measured using the coin sample on a laser flash system (Netzsch LFA 457, Germany). The specific heat capacity (Cp) was measured on a differential scanning calorimetry thermal analyzer (Netzsch DSC 404 C, Germany). The relative density, determined by the Archimedes method, were higher than 97% for all the samples. The uncertainty for the electrical conductivity is 3%, the Seebeck coefficient 5%, the thermal conductivity 7%, so the combined uncertainty for the power factor is 13% and that for ZT value is 20%. To increase the readability of the curves, error bars were not shown in the figures. The roomtemperature and high-temperature Hall Coefficient RH was measured by using the PPMS (Physical Properties Measurement System, Quantum Design) and home-made instrument, respectively. The Hall carrier concentration (nH) was obtained by nH ¼ 1/ eRH and the Hall carrier mobility (mH) was calculated by s ¼ emHnH, where e is the electronic charge and s the electrical conductivity. 3. Results Fig. 1 presents the X-ray diffraction (XRD) patterns of hotpressed Zr3Ni3Sb4-xTex disk samples (x ¼ 0, 0.025, 0.05, 0.1, and 0.2). The diffraction peaks of all the samples can be well indexed to Y3Au3Sb4-type structure without any impurity phase within the detection limit of the XRD spectrometer (Fig. 1a) [53]. Due to the fact that the ionic radius of Te (221 Å) is smaller than that of Sb (245 Å), the calculated lattice parameter a gradually decreases with increasing Te doping concentration up to 0.1 (Fig. 1b), consistent with the Vegard's law. However, this saturation behavior between x ¼ 0.1 and x ¼ 0.2 samples may be related with the finite solubility of Te on the Sb site, which will also influence the thermoelectric properties that will be discussed later. The typical fractured surface morphology of Zr3Ni3Sb3.95Te0.05 sample shows the characteristic of intergranular brittle fracture without any observable micro-hole (Fig. 2a). The grain size ranges from 500 nm to 3 mm, similar with other nanostructured medium and high thermoelectric materials [20,56]. The corresponding elemental mapping image indicates that all the elements (Zr, Ni, Sb, and Te) are distributed homogeneously without obvious element aggregation. Fig. 3 shows the thermoelectric properties of Zr3Ni3Sb4-xTex samples. It is obvious that undoped Zr3Ni3Sb4 sample is an intrinsic p-type semiconductor because of its positive Seebeck coefficients at the whole measured temperature range (Fig. 3a), consistent with

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Fig. 1. (a) XRD patterns and (b) enlarged XRD patterns from 76 to 80 for Zr3Ni3Sb4samples (x ¼ 0, 0.025, 0.05, 0.1 and 0.2).

xTex

the majority of Zintl antimonides [57]. Besides, the Seebeck coefficient of undoped Zr3Ni3Sb4 monotonically decreases with the increased temperature. In contrast, all the Te doped Zr3Ni3Sb4 samples show the n-type conduction behavior with negative Seebeck coefficients, even for the lowest Te doping (x ¼ 0.025), which indicates that Te is a strong n-type dopant in the Zr3Ni3Sb4 system. With increasing Te doping concentration from x ¼ 0.025 to x ¼ 0.2, the Seebeck coefficient gradually decreases owing to the significantly increased carrier concentration from ~5  1018 cm3 to ~3.7  1020 cm3. In addition, high Te doping concentration, namely high carrier concentration, could obviously suppress the bipolar effect. Simultaneously, the accompanied ionized impurity scattering results in decreased Hall carrier mobility from ~17.2 cm2 V1 s1 to ~4.9 cm2 V1 s1. As shown in Fig. 3b, both undoped Zr3Ni3Sb4 and Zr3Ni3Sb3.975Te0.025 sample exhibit a typical semiconductor behavior, while for the other samples the electrical resistivity first decreases and then almost keeps constant with increasing the temperature. The different temperature dependent electrical resistivity indicates the changed carrier scattering

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mechanism that will be thoroughly discussed later. With increasing Te doping concentration from x ¼ 0.025 to x ¼ 0.1, the electrical resistivity monotonically decreases because of the sharply increased carrier concentration. However, the electrical resistivity of Zr3Ni3Sb3.8Te0.2 sample shows the abnormal increment in comparison with that of Zr3Ni3Sb3.9Te0.1 sample, which may be caused by the appearance of minor impurity phases beyond XRD detection. Consequently, power factor shows a significant enhancement upon Te doping due to the optimized carrier concentration, especially in the high temperature range (Fig. 3c). For example, power factor at 773 K of Zr3Ni3Sb3.95Te0.05 sample (~17.2 mW cm1 K2) is ~300 times higher than that of undoped Zr3Ni3Sb4 sample (~0.05 mW cm1 K2). It should be pointed out that both n-type and p-type Zr3Ni3Sb4 based materials exhibit a relatively higher power factor compared with other Zintl antimonides [4], where the maximum power factor of Te doping (n-type), Co doping (p-type) and Cu doping (n-type) is around ~18.7 mW cm1 K2, ~20.5 mW cm1 K2, and ~17.6 mW cm1 K2 [51,52], respectively. This can be ascribed to the high symmetry of Zr3Ni3Sb4 crystal structure (cubic) while most of Zintl antimonides have a low crystal symmetry (orthorhombic or tetragonal structure) [40]. In terms of practical application, higher power factor leads to larger output power density of device at given working boundary conditions [13]. Thus, Zr3Ni3Sb4 based materials are highly applicable for intermediate-temperature power generation, especially when the heat source is unlimited (such as solar heat), or heat source is free (such as waste heat from automobiles, steel industry, etc.). Due to the complex crystal structure and large unit cell of Zr3Ni3Sb4 in comparison with half-Heusler alloys, e.g. ZrNiSn shown as dash line [58,59], much lower total thermal conductivity can be obtained (Fig. 3d). Upon Te doping, the increased total thermal conductivity mainly stems from the increased electronic thermal conductivity. As shown in Fig. 3e, Te doping obviously enhances the ZT and the highest peak ZT at 773 K is about 0.6 for Zr3Ni3Sb3.95Te0.05 sample. In practice, the conversion efficiency of power generation is dominated by the average ZT over the imposed temperature difference. The average ZT from 300 K to 773 K of Zr3Ni3Sb3.95Te0.05 sample (~0.33) is almost 5 times higher than that of Zr3Ni3Sb4 sample (~0.07) (Fig. 3f). Further enhancement of ZT may be highly possible via isoelectronic alloying to largely reduce thermal conductivity, e.g., Hf alloying on the Zr site or Pt alloying on Ni site [52,54].

4. Discussion Carrier concentration plays a critical role in the electrical transport properties of thermoelectric materials [17], so the room-

Fig. 2. (a) Typical SEM image of the fracture morphology and (b) the corresponding EDS mapping for Zr3Ni3Sb3.95Te0.05 sample.

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Fig. 3. Temperature dependent (a) Seebeck coefficient, (b) electrical resistivity, (c) power factor, (d) total thermal conductivity and (e) ZT; (f) average ZT for Zr3Ni3Sb4-xTex samples.

temperature carrier concentration dependence of doping concentration for both Te and Cu doping is shown in Fig. 4 [51]. It is evident that the doping effectiveness of Te is significantly higher than that of Cu. For example, the carrier concentration of Zr3Ni3Sb3.9Te0.1 (~2.5  1020 cm3) is almost an order of magnitude higher than that of Zr3Ni2.7Cu0.3Sb4 (~2.6  1019 cm3) [51]. This indicates that Te behaves as a much stronger n-type donor in Zr3Ni3Sb4 than Cu. Tamaki et al. claimed that Cu forms the deep donor level and then a fraction of doping induced electrons can be thermally excited into the conduction band [51], which could explain the lower carrier concentration at Cu doping condition. It should also be mentioned

that Te has a lower solid solubility than Cu in Zr3Ni3Sb4. Fig. 5 compares the electrical transport properties of Zr3Ni3Sb3.95Te0.05 and Zr3Ni2.7Cu0.3Sb4 samples [51]. For the temperature dependent Seebeck coefficient (Fig. 5a), both samples show a similar decreased trend at the low temperature range (from 300 K to 573 K). At the elevated temperature (from 573 K to 773 K), however, the absolute value of Seebeck coefficient of Zr3Ni2.7Cu0.3Sb4 sample obviously decreases in turn, in sharp contrast to the almost linearly increased trend for Zr3Ni3Sb3.95Te0.05 sample. It means that Zr3Ni2.7Cu0.3Sb4 sample exhibits a stronger bipolar effect, compared with Zr3Ni3Sb3.95Te0.05 sample. Generally, bipolar effect can be frequently observed at ordinary temperatures for narrow-gap semiconductors or at high temperatures for wide-gap semiconductors, where two types of carriers, the majority and the minority, contribute to the current flow [12]. Thus, the actual Seebeck coefficient becomes a weighted average of the Seebeck coefficients related to two types of carriers, shown in the following Equation (1):

a¼

Fig. 4. Doping concentration dependent carrier concentration for both Te and Cu doping (Tamaki et al.) [50].

Sn sn þ Sp sp sn þ sp

(1)

where sn and Sn (sp and Sp) are the partial electrical conductivity and Seebeck coefficient for electrons (holes), respectively [12]. Since Seebeck coefficient of electrons and holes have opposite sign, the concurrent flow of both carriers would enable the actual Seebeck effect cancel each other, leading to the signature of maximum

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Fig. 5. Temperature dependent electrical transport properties for Zr3Ni3Sb3.95Te0.05 and Zr3Ni2.7Cu0.3Sb4 (Tamaki et al.) [50]. (a) Seebeck coefficient, (b) electrical resistivity, (c) carrier mobility and (d) power factor.

Seebeck coefficient corresponding to a finite temperature. Here at the intrinsic conduction regime, the Goldsmid-Sharp formula, Eg ¼ 2ejSjmaxTmax, is employed to estimate the band gap Eg from measured high-temperature Seebeck coefficient, where Tmax is the measured temperature at which the maximum of the absolute thermopower (jSjmax) occurs, e is the charge of electron. The obtained band gap of Zr3Ni3Sb3.95Te0.05 is about 0.28 eV, much larger than that of undoped (0.22 eV) and Zr3Ni2.7Cu0.3Sb4 (0.19 eV) samples. This should be related with Burstein-Moss effect [60]. In the case of heavily doped n-type semiconductors, like our Te doped Zr3Ni3Sb4, Fermi level can be located in conduction band Ec þ dE (strong degenerated semiconductors). Thus, the optical transition of electrons from valance band can take place only to the energy levels in conduction band located above the Fermi level, since energy levels below the Fermi level are filled with electrons. In this case, the location of the self-absorption edge shifts toward higher energy and thus the effective band gap is equal to Eg þ dE. But this explanation requires further first-principle calculations to support that is beyond our current study. For undoped Zr3Ni3Sb4, the narrow band gap and the small weighted mobility ratio between majority and minority give rise to a strong bipolar effect [52,54], resulting in the gradually decreased Seebeck coefficient at the entire measured temperature range. Even though Cu doping could weaken the bipolar effect to a certain extent, it is still apparent, particularly at the high temperature. For Te doping, however, the bipolar effect almost vanishes due to the combination of higher carrier concentration and larger band gap. As shown in Fig. 5b, Zr3Ni3Sb3.95Te0.05 and Zr3Ni2.7Cu0.3Sb4 show a significantly different temperature dependence of electrical resistivity from 300 K to 050 K while the high-temperature behaviors are nearly identical from 500 K to 773 K. As expected, the measured hightemperature carrier mobility of Zr3Ni3Sb3.95Te0.05 shows a typical temperature exponent g of 1.5, indicative of ionized impurity scattering of the carriers, while the temperature exponent g of 0 for Zr3Ni2.7Cu0.3Sb4 demonstrates the mixed acoustic phonon and ionized impurity scattering at the low temperature range (Fig. 5b and c). The underlying reason of diverse carrier scattering mechanism for different elements doping is still an open question that is

beyond the present study, partly resulted from the higher actual ionized scatters for Te doping. At the high temperature range, both samples show an identical temperature exponent g of 0.5, still as the mixed acoustic phonon and ionized impurity scattering but acoustic phonon scattering works stronger. Finally, Zr3Ni2.7Cu0.3Sb4 shows a higher power factor at the low temperature range, while Zr3Ni3Sb3.95Te0.05 shows a higher power factor at the high temperature range (Fig. 5d). To investigate the carrier scattering mechanism and carrier concentration on the electrical transport properties of Zr3Ni3Sb4based materials [51], Fig. 6 presents the carrier mobility and power factor depending on scattering exponent of electrical resistivity or carrier concentration. Intuitively, Te doping leads to a much lower carrier mobility at 300 K, as a consequence of strong ionized impurity scattering, compared with that of Cu doping with mixed acoustic phonon and ionized impurity scattering (Fig. 6a). Normally, ionized impurity scattering dramatically deteriorates the carrier mobility in semiconductors, where the carrier scattering probability is inversely proportional to E3/2 (E: carrier energy) or v3 (v: drift velocity) [21,61]. Normally, acoustic phonon scattering is the dominating factor for carrier in thermoelectrics, e.g. PbX (X ¼ Te, Se or S) system, which is caused by the low dopant concentration and small dielectric constant [62]. For the specific carrier scattering mechanism, the carrier mobility is also closely related to the carrier concentration at a certain temperature (Fig. 6b). Due to the appearance of ionized impurity centers, Te doping has a relatively higher room-temperature Seebeck coefficient at the same carrier concentration (Fig. 6c). In addition, the high-temperature Seebeck coefficient of Te doping is obviously higher than Cu doping because of the weaker bipolar effect (Fig. 6c). Collectively, since the sharply decreased carrier mobility prevails over the increased Seebeck coefficient at room temperature, Te doping has a lower power factor (Fig. 6d). Our study clearly points out that ionized impurity scattering to charge carrier is unfavorable for the electrical transport properties, consistent with recent work of ntype Zintl Mg3Sb2-based materials [23]. Despite the same hightemperature carrier scattering mechanism, Te doping, however, leads to the higher power factor, which can be mainly ascribed to

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Fig. 6. The role of carrier scattering mechanism on the electrical transport properties for Te and Cu doping (Tamaki et al.) [50] (a) Relationship between temperature dependent scattering exponential of electrical resistivity and carrier mobility, (b) relationship between carrier concentration and carrier mobility, (c) relationship between carrier concentration and absolute value of Seebeck coefficient, (d) relationship between temperature dependent scattering exponential of electrical resistivity and power factor.

the weaker bipolar effect (Fig. 6d). Fig. 7 compares the different role of Te doping and Cu doping in thermal transport properties [51]. ktot of both undoped and Cu doped samples show an obvious uptrend at the relatively high temperature range while this phenomenon is almost diminished for Te doping, which indicate that Te doping could strongly suppress or eliminate the bipolar effect. As known, kele is proportional

to the electrical conductivity (s) through the Wiedemann-Franz relationship, kele ¼ LsT, where L is the Lorenz number. Since the carrier scattering mechanism is rather complex for Zr3Ni3Sb4 system, it is difficult to calculate the L using the theoretical model. Thus, 1.5  108 W U K2 (for non-degenerate semiconductors) and L ¼ 2.0  108 W U K2 (for degenerate semiconductors) are employed in the Wiedemann-Franz formula for undoped Zr3Ni3Sb4

Fig. 7. Comparison of temperature dependent (a) total thermal conductivity, (b) lattice thermal conductivity, (c) bipolar thermal conductivity and (d) ZT for Zr3Ni3Sb4, Zr3Ni3Sb3.95Te0.05 and Zr3Ni2.7Cu0.3Sb4 alloys (Tamaki et al.) [50].

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and doped Zr3Ni3Sb4, respectively. The method proposed by Kitagawa et al. is utilized to roughly separate the contribution of klat and kbip [31,63]. Umklapp scattering, as the predominant mechanism for phonons above Debye temperature for most thermoelectric materials, determines the relationship between the klat and the reciprocal temperature T1 being linear. Provided that bipolar effect does not exist at the low temperature range, high-temperature klat and kbip can be roughly estimated by extrapolating this linear relationship between klat and T1. Bipolar thermal conductivity in semiconductors can be expressed as

kbip ¼


2 sn sp S  Sn T sn þ sp p

(2)

Therefore, the existence of majority and minority significantly strengthens heat conduction. In the intrinsic regime, kbip should increase exponentially with increasing temperature, shown in Equation (3):



kbip ¼ exp 

Eg 2kB T

 (3)

where kB is the Boltzmann constant. That is why kbip of undoped Zr3Ni3Sb4 dramatically increases at the elevated temperature. Evidently, a significantly low kbip is observed for Te doping, in contrast with the high kbip of Cu doping (Fig. 7c). For instance, the peak kbip of Zr3Ni3Sb3.95Te0.05 (~0.2 W m1 K1) is one-third of Zr3Ni2.7Cu0.3Sb4 (~0.6 W m1 K1) and nearly one-fifth of undoped one (~0.9 W m1 K1). This critical suppression should be ascribed to the combination of high majority concentration and enlarged band gap upon Te doping. Therefore, Te doping resulted in a higher ZT at the high temperature range (Fig. 7d), caused by the weakening or vanishing of bipolar effect. 5. Conclusion In summary, nanostructured Te-doped Zintl Zr3Ni3Sb4 alloys were directly fabricated by mechanical alloying and hot pressing. Te doping could convert the intrinsic p-type characteristic to n-type. The different carrier scattering mechanisms on the lowtemperature electrical transport properties were also demonstrated, where both carrier mobility and power factor by Te doping, due to the detrimental effect of ionized scattering, are lower than that of Cu doping in which the mixed acoustic phonon and ionized impurity scattering dominates. More significantly, both high majority-carrier concentration and enlarged band gap lead to the significant suppression of high-temperature bipolar effect upon Te doping. Collectively, the relatively high ZT of ~0.6 at 773 K can be obtained for Zr3Ni3Sb3.95Te0.05, almost double that of Zr3Ni2.7Cu0.3Sb4. Acknowledgement This work was supported by “Solid State Solar Thermal Energy Conversion Center (S3TEC)”, an Energy Frontier Research Center funded by the U.S. Department of Energy, Office of Science, Office of Basic Energy Science under award number DE-SC0001299. References [1] M.S. Dresselhaus, G. Chen, M.Y. Tang, R.G. Yang, H. Lee, D.Z. Wang, Z.F. Ren, J.P. Fleurial, P. Gogna, Adv. Mater. 19 (2007) 1043e1053. [2] G.J. Snyder, E.S. Toberer, Nat. Mater. 7 (2008) 105e114. [3] J. Mao, Z.H. Liu, Z.F. Ren, npj, Quantum Mater. 1 (2016) 16028. [4] J. Shuai, J. Mao, S.W. Song, Q.Y. Zhang, G. Chen, Z.F. Ren, Mater. Today. Phys. 1 (2017) 74e95. [5] W.S. Liu, J.Z. Hu, S.M. Zhang, M.J. Deng, C.G. Han, Y. Liu, Mater. Today. Phys. 1

7

(2017) 50e60. [6] B. Poudel, Q. Hao, Y. Ma, Y.C. Lan, A. Minnich, B. Yu, X. Yan, D.Z. Wang, A. Muto, D. Vashaee, X.Y. Chen, J.M. Liu, M.S. Dresselhaus, G. Chen, Z.F. Ren, Science 320 (2008) 634e638. [7] D. Kraemer, J.H. Sui, K. McEnaney, H.Z. Zhao, Q. Jie, Z.F. Ren, G. Chen, Energy Environ. Sci. 8 (2015) 1299e1308. [8] Z.H. Liu, H.Y. Geng, J. Mao, S. Jing, R. He, C. Wang, W. Cai, J.H. Sui, Z.F. Ren, J. Mater. Chem. A 4 (2016) 16834e16840. [9] L.D. Zhao, G.J. Tan, S.Q. Hao, J.Q. He, Y.L. Pei, H. Chi, H. Wang, S.K. Gong, H.B. Xu, V.P. Dravid, C. Uher, G.J. Snyder, C. Wolverton, M.G. Kanatzidis, Science 351 (2016) 141e144. [10] Z.W. Chen, Z.Z. Jian, W. Li, Y.J. Chang, B.H. Ge, R. Hanus, J. Yang, Y. Chen, M.X. Huang, G.J. Snyder, Y.Z. Pei, Adv. Mater. 29 (2017), 1606768-n/a. [11] X.F. Meng, Z.H. Liu, B. Cui, D.D. Qin, H.Y. Geng, W. Cai, L.W. Fu, J.Q. He, Z.F. Ren, J.H. Sui, Adv. Energy Mater. 7 (2017), 1602582. [12] H.J. Goldsmid, Introduction to Thermoelectricity, Springer Science & Business Media, New York, 2009. [13] W.S. Liu, H.S. Kim, Q. Jie, Z.F. Ren, Scr. Mater. 111 (2016) 3e9. [14] Z.H. Liu, Y.M. Wang, J. Mao, H.Y. Geng, J. Shuai, Y.X. Wang, R. He, W. Cai, J.H. Sui, Z.F. Ren, Adv. Energy Mater. 6 (2016), 1502269. [15] T.J. Zhu, Y.T. Liu, C.G. Fu, J.P. Heremans, J.G. Snyder, X.B. Zhao, Adv. Mater. 29 (2017), 1605884. [16] Y.Z. Pei, A. LaLonde, S. Iwanaga, G.J. Snyder, Energy Environ. Sci. 4 (2011) 2085e2089. [17] Y.Z. Pei, H. Wang, G.J. Snyder, Adv. Mater. 24 (2012) 6125e6135. [18] Y. Pei, Z.M. Gibbs, A. Gloskovskii, B. Balke, W.G. Zeier, G.J. Snyder, Adv. Energy Mater. 4 (2014) 9201e9210. [19] Q. Zhang, E.K. Chere, K. McEnaney, M.L. Yao, F. Cao, Y.Z. Ni, S. Chen, C. Opeil, G. Chen, Z.F. Ren, Adv. Energy Mater. 5 (2015), 1401977. [20] R. He, H.T. Zhu, J.Y. Sun, J. Mao, H. Reith, S. Chen, G. Schierning, K. Nielsch, Z.F. Ren, Mater. Today. Phys. 1 (2017) 24e30. [21] S.Y. Wang, J. Yang, L.H. Wu, P. Wei, W.Q. Zhang, J.H. Yang, Adv. Funct. Mater. 25 (2015) 6660e6670. [22] L. Pan, S. Mitra, L.D. Zhao, Y.W. Shen, Y.F. Wang, C. Felser, D. Berardan, Adv. Funct. Mater. 26 (2016) 5149e5157. [23] J. Shuai, J. Mao, S.W. Song, Q. Zhu, J.F. Sun, Y.M. Wang, R. He, J.W. Zhou, G. Chen, D.J. Singh, Z.F. Ren, Energy Environ. Sci. 10 (2017) 799e807. [24] K.F. Hsu, S. Loo, F. Guo, W. Chen, J.S. Dyck, C. Uher, T. Hogan, E.K. Polychroniadis, M.G. Kanatzidis, Science 303 (2004) 818e821. [25] X. Yan, W.S. Liu, H. Wang, S. Chen, J. Shiomi, K. Esfarjani, H.Z. Wang, D.Z. Wang, G. Chen, Z.F. Ren, Energy Environ. Sci. 5 (2012) 7543e7548. [26] Z.H. Liu, Y.S. Zhang, J. Mao, W.H. Gao, Y.M. Wang, J. Shuai, W. Cai, J.H. Sui, Z.F. Ren, Acta Mater. 128 (2017) 227e234. [27] J. Mao, Y.M. Wang, B.H. Ge, Q. Jie, Z.H. Liu, U. Saparamadu, W.S. Liu, Z.F. Ren, Phys. Chem. Chem. Phys. 18 (2016) 20726e20737. [28] K.P. Zhao, P.F. Qiu, Q.F. Song, A.B. Blichfeld, E. Eikeland, D. Ren, B.H. Ge, B.B. Iversen, X. Shi, L.D. Chen, Mater. Today. Phys. 1 (2017) 14e23. [29] Z.W. Chen, B.H. Ge, W. Li, S.Q. Lin, J.W. Shen, Y.J. Chang, R. Hanus, G.J. Snyder, Y.Z. Pei, Nat. Commun. 8 (2017) 13828. [30] H. Wang, Y.Z. Pei, A.D. LaLonde, G.J. Snyder, Adv. Mater. 23 (2011) 1366e1370. [31] L.D. Zhao, H.J. Wu, S.Q. Hao, C.I. Wu, X.Y. Zhou, K. Biswas, J.Q. He, T.P. Hogan, C. Uher, C. Wolverton, Energy Environ. Sci. 6 (2013) 3346e3355. [32] Z.H. Liu, H.Y. Geng, J. Shuai, Z.Y. Wang, J. Mao, D.Z. Wang, Q. Jie, W. Cai, J.H. Sui, Z.F. Ren, J. Mater. Chem. C 3 (2015) 10442e10450. [33] H.R. Yang, J.-H. Bahk, T. Day, A.M. Mohammed, G.J. Snyder, A. Shakouri, Y. Wu, Nano Lett. 15 (2015) 1349e1355. [34] L.B. Zhang, P.H. Xiao, L. Shi, G. Henkelman, J.B. Goodenough, J.S. Zhou, J. Appl. Phys. 117 (2015), 155103. €hls, U. Aydemir, P.F. Qiu, C.C. Stoumpos, R. Hanus, [35] S.D. Kang, J.-H. Po M.A. White, X. Shi, L.D. Chen, M.G. Kanatzidis, G.J. Snyder, Mater. Today. Phys. 1 (2017) 7e13. [36] S.Y. Wang, J. Yang, T. Toll, J.H. Yang, W.Q. Zhang, X.F. Tang, Sci. Rep. 5 (2015). [37] D.M. Rowe, CRC Handbook of Thermoelectrics, CRC press, New York, 1995. [38] G. Nolas, D. Morelli, T.M. Tritt, Annu. Rev. Mater. Sci. 29 (1999) 89e116. [39] G.J. Snyder, M. Christensen, E. Nishibori, T. Caillat, B.B. Iversen, Nat. Mater. 3 (2004) 458e463. [40] E.S. Toberer, A.F. May, G.J. Snyder, Chem. Mater. 22 (2009) 624e634. [41] T. Takabatake, K. Suekuni, T. Nakayama, E. Kaneshita, Rev. Mod. Phys. 86 (2014) 669e716. [42] S.M. Kauzlarich, S.R. Brown, G.J. Snyder, Dalton Trans. (2007) 2099e2107. [43] R. Nesper, Z. Anorg. Allg. Chem. 640 (2014) 2639e2648. [44] S.R. Brown, S.M. Kauzlarich, F. Gascoin, G.J. Snyder, Chem. Mater. 18 (2006) 1873e1877. [45] F. Gascoin, S. Ottensmann, D. Stark, S.M. Haïle, G.J. Snyder, Adv. Funct. Mater. 15 (2005) 1860e1864. [46] E.S. Toberer, A.F. May, B.C. Melot, E. Flage-Larsen, G.J. Snyder, Dalton Trans. 39 (2010) 1046e1054. [47] A. Zevalkink, E.S. Toberer, W.G. Zeier, E. Flage-Larsen, G.J. Snyder, Energy Environ. Sci. 4 (2011) 510e518. [48] G.S. Pomrehn, A. Zevalkink, W.G. Zeier, A. van de Walle, G.J. Snyder, Angew. Chem. Int. Ed. 53 (2014) 3422e3426. [49] L. Bjerg, G.K. Madsen, B.B. Iversen, Chem. Mater. 23 (2011) 3907e3914. [50] H. Tamaki, H.K. Sato, T. Kanno, Adv. Mater. 28 (2016) 10182e10187. [51] H. Tamaki, T. Kanno, A. Sakai, K. Takahashi, H. Kusada, Y. Yamada, Appl. Phys. Lett. 104 (2014), 122103.

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Z. Liu et al. / Materials Today Physics xxx (2017) 1e8

[52] H. Tamaki, T. Kanno, A. Sakai, K. Takahashi, Y. Yamada, J. Appl. Phys. 118 (2015), 055103. [53] M. Wang, R. McDonald, A. Mar, Inorg. Chem. 38 (1999) 3435e3438. [54] P. Larson, S. Mahanti, J. Salvador, M. Kanatzidis, Phys. Rev. B 74 (2006), 035111. [55] J. Salvador, X. Shi, J. Yang, H. Wang, Phys. Rev. B 77 (2008), 235217. [56] Q. Zhang, B.L. Liao, Y.C. Lan, K. Lukas, W.S. Liu, K. Esfarjani, C. Opeil, D. Broido, G. Chen, Z.F. Ren, Proc. Natl. Acad. Sci. U. S. A. 110 (2013) 13261e13266. [57] L. Bjerg, G.K. Madsen, B.B. Iversen, Chem. Mater. 24 (2012) 2111e2116.

[58] W.G. Zeier, J. Schmitt, G. Hautier, U. Aydemir, Z.M. Gibbs, C. Felser, G.J. Snyder, Nat. Rev. Mater. (2016) 16032. [59] T.J. Zhu, C.G. Fu, H.H. Xie, Y.T. Liu, X.B. Zhao, Adv. Energy Mater. 5 (2015), 1500588. [60] S.M. Sze, K.K. Ng, Physics of Semiconductor Devices, John wiley & sons, 2006. [61] P. Debye, E. Conwell, Phys. Rev. 93 (1954) 693. [62] D. Khokhlov, Lead chalcogenides: Physics and Applications, CRC Press, 2002. [63] H. Kitagawa, M. Wakatsuki, H. Nagaoka, H. Noguchi, Y. Isoda, K. Hasezaki, Y. Noda, J. Phys. Chem. Solids 66 (2005) 1635e1639.

Please cite this article in press as: Z. Liu, et al., Tellurium doped n-type Zintl Zr3Ni3Sb4 thermoelectric materials: Balance between carrierscattering mechanism and bipolar effect, Materials Today Physics (2017), https://doi.org/10.1016/j.mtphys.2017.08.002