Method of Efficient Ag Doping for Fermi Level Tuning of

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The Journal of Physical Chemistry

Method of Efficient Ag-Doping for Fermi Level Tuning of Thermoelectric Bi0.5Sb1.5Te3 Alloys using a Chemical Displacement Reaction Sungho Seo1, Kyungseok Lee2, Youngkeun Jeong2, Min-Wook Oh3*, Bongyoung Yoo1,4* 1

Department of Bionanotechnology, Hanyang University, Ansan, Gyeonggi-do, Republic of

Korea 2

National Core Research Center for Hybrid Materials Solution, Pusan National University,

Pusan, Republic of Korea 3

Department of Advanced Materials Engineering, Hanbat National University, Daejeon,

Republic of Korea 4

Department of Materials Engineering, Hanyang University, Ansan, Gyeonggi-do, Republic of

Korea *Co-corresponding authors [email protected] (Dr. Bongyoung Yoo), T: 82-31-400-5229 [email protected] (Dr. Min-Wook Oh)

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Keyword: thermoelectric, Bi0.5Sb1.5Te3, Ag doping, displacement reaction, Fermi level

Abstract

Silver is a well-known element for use as a p-type dopant in Bi2Te3-related compounds. In this paper, an efficient method for incorporating ultra-low Ag dopant concentrations (< 1,300 ppm) into Bi0.5Sb1.5Te3 via a simple chemical displacement reaction is described. Powders of Bi0.5Sb1.5Te3 synthesized by mechanical alloying were reacted with Ag+ ions in dilute HNO3 solutions (pH 0.2), resulting in the deposition of Ag on the surface of the powders due to the difference in reduction potential between Ag and Bi0.5Sb1.5Te3. The Ag/Bi0.5Sb1.5Te3 powders were then sintered by SPS, and the thermoelectric properties of the dense Ag-doped samples were measured. Low Ag-doped samples showed behavior characteristic of partially degenerate semiconductors, while highly-doped specimens exhibited properties associated with fully degenerate semiconductors. From the measured transport properties and theoretical estimations, successful tuning of the Fermi level with Ag was confirmed. Consequently, the temperature at which the peak dimensionless figure of merit (ZT) value was obtained increased from 50 °C to 250 °C. Such findings may be beneficial in the utilization of waste heat over a wide temperature range, as the Ag-doped samples could be employed as functionally graded materials for thermoelectric modules.


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Introduction Thermoelectrics are a promising class of materials for the realization of power generation via waste heat recovery and refrigeration without using environmentally harmful refrigerants.1-3 The efficiency of thermoelectrics is represented by the dimensionless figure of merit (ZT), defined as ZT = (σ·S2/κ)T, where σ, S, κ, and T are the electrical conductivity, Seebeck coefficient, thermal conductivity, and average temperature, respectively.4-5 The figure of merit can be also represented in terms of the Fermi energy using classical statistics as follows:6 =

·

=

&=

∗ "

#

$%

(Eq. 1)

,

$ ' " ( % # )

,

where λ is a scattering parameter (0 for acoustic phonon scattering), η is the reduced Fermi energy (=Ef/kBT), kB is the Boltzmann constant, µ is the carrier mobility, κl is the lattice thermal conductivity, m* is the effective carrier mass, m0 is the electron mass, and h is Planck’s constant. As shown in Eq. 1, tuning of the Fermi energy is necessary to achieve high ZT. In this study, Bi0.5Sb1.5Te3 (BST) and Ag were selected as the thermoelectric material and dopant for tuning the Fermi energy, respectively. BST is a representative thermoelectric for room temperature applications7, while Ag is a well-known p-type dopant for BST. I. Klichova et al.


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and J. K. Lee et al. attempted to synthesize Ag-doped BST by melting all of the constituent elements.8-9 K. H. Lee et al. decorated BST powders with Ag nanoparticles via thermal decomposition; the resulting powders were then sintered.10 Here, BST was coated with Ag by a simple chemical displacement reaction. In particular, Ag+ ions were reduced by a displacement reaction with BST, thereby facilitating the deposition of Ag on the surface of the BST powders (Fig. 1(a)). The as-synthesized powders were subsequently sintered to produce Ag-doped BST specimens. Compared to techniques described in previous studies, the method introduced here has two advantages. Firstly, it is applicable to Bi0.5Sb1.5Te3 having various shape such as nanoplates, nanorods and nanowires (and also other novel shape), synthesized by other researchers. Secondly, the strategy allows for efficient control of the BST carrier concentration down to very low Ag amounts. As shown in Fig. 1(b), the rate of change in the hole concentration with respect to the Ag content (=average slope) was 4 ~8 times larger in this work than in previous studies, which means that the doping efficiency was much higher by comparison.

Experimental Section Ternary BST powders were synthesized by mechanical alloying from Bi, Sb, and Te elemental chunks with high purity (4~5N).11 The elemental chunks were milled together by high-energy ball milling (attrition mill) for 24 h in an Ar atmosphere. The phase of the milled powders was consistent with that of stoichiometric Bi0.5Sb1.5Te3.0 alloys (Fig. S1(a)), while the particle size ranged from hundreds of nanometers to a few micrometers, as shown in Fig. S1(b).


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In order to utilize the displacement reaction and deposit Ag on the surface of the BST powders, dilute HNO3 (adjusted to have a pH of 0.2 so as to prevent oxidation of the BST powders) and AgNO3 (dissolved in the dilute HNO3) solutions were prepared. It should be noted that surface oxidation of the BST was not thoroughly avoided even at a solution pH of 0.2, as a small decrease in the electrical conductivity of BST was observed after an HNO3 solution treatment with no AgNO3 addition. The HNO3 solutions (800 mL) in a 1,000 mL beaker were mechanically stirred (300 rpm) at room temperature, and then BST powders (20.0 g) and AgNO3 solutions were added. The bath was sealed with paraffin film to prevent evaporation of the solution during the reaction. After allowing the solution contents to react for 12 h, the BST powders were collected by centrifugation. Residual moisture was then removed by a vaporization process carried out with a vacuum pump. The powders were sintered by spark plasma sintering (SPS) at 380 °C for 10 min using an SPS-825 apparatus (Fuji Electronic Industrial Co., Ltd.) so as to yield bulk cylindrical samples (12.7 mm in diameter and 20 mm in thickness); the heating rate and pressure were 50 °C/min and 40 MPa, respectively. The bulk specimens were machined into bar-shaped pieces (3 × 3 × 10 mm3) for measurement of the electrical conductivity and Seebeck coefficient, as well as a disk-shaped pieces (12.7 mm in diameter and 2 mm in thickness) for thermal diffusivity tests. The phase of powders was analyzed by X-ray diffractometry (XRD, B/max 2500, Rigaku). Low-magnification images of the powders and fractured samples were obtained with a field effect scanning electron microscope (FE-SEM, MIRA3, Tescan). Inductively coupled plasma atomic emission spectroscopy (ICP-AES, Spectro Arcos, Spectro) analysis was carried out to measure the ultra-low Ag concentration. Both the electrical conductivity and Seebeck coefficient of the specimens were measured with a ZEM-3 (ULVAC-RIKO) apparatus. The thermal


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conductivity was calculated from the equation κ = aCPd (a: thermal diffusivity, Cp: specific heat capacity, d: sample density). The thermal diffusivity was experimentally measured by the laser flash method (LFA-447, Netzsch), while the specific heat capacity was theoretically obtained according to computations of the Cp values for Bi2Te3 and Sb2Te3.12 Densities were calculated based on the specimen geometry, and all samples exhibited values in the range of 95.1 ± 1.3% of the theoretical density. Carrier concentration and mobility were measured by a Hall effect measurement system (HMS-3000, Ecopia) equipped with a 0.55 T electromagnet. Highmagnification images acquired with a scanning transmission electron microscope (STEM, JEM2100F, JEOL) were used to examine the microstructure of sintered specimens.

Results and Discussion The aforementioned displacement reaction occurs due to the difference in reduction potential between the Ag+ ion and BST powder (EAg = + 0.7996 V, EBST = - 0.038 ~ + 0.142 V (vs. NHE)).13-14 A lower reduction potential means a higher ionization energy and thus, BST tends to be ionized by donating electrons to Ag+ ions, i.e., Ag+ ions are reduced on the surface of BST. In order to confirm the reduction of Ag on the surface of BST, the Ag distribution in the sample prepared with 6 mmol of AgNO3 was investigated by EDS mapping. As shown in Fig. 2, the BST surface was homogeneously coated with Ag, the concentration of which was determined to be ~30,000 ppm. Furthermore, no secondary or elemental Ag phases were observed in the XRD pattern (Fig. S2). Such a finding may be attributed to the low Ag concentration and thin, homogeneous coating of Ag on the surface of the BST powders. The added and detected Ag concentrations are listed in Table 1. The detected Ag concentration linearly increased with the


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addition of AgNO3. In the case of sample #2, the detected value for the Ag concentration was not viewed as reliable due to the detection limits associated with the ICP-AES system. Fig. 3 shows the electrical conductivity of the samples with respect to temperature. As the temperature increased, the electrical conductivity of all samples decreased and, with the exception of samples #4 and #5, tended to saturate at high temperature. The observed decrease in the electrical conductivity with temperature is typical behavior of degenerate semiconductors. Over the entire range of testing temperatures, the electrical conductivity increased at higher Ag concentrations. The electrical conductivity of a p-type semiconductor is proportional to the hole carrier concentration (nh) and the mobility (µh) according to σ = e·nh·µh, where e is the elementary charge. Therefore, the carrier concentration and mobility were estimated via Hall coefficient (RH) measurements. The carrier concentration was computed using the expression nh=1/qRH and assuming single and parabolic bands, as is common for Bi2Te3-related compounds.15-16 With an increase in the Ag content, the carrier concentration also increased while the mobility exhibited a decrease (Fig. S3). Thus, the increase in the electrical conductivity is mainly attributed to an increase in carrier concentration with the addition of Ag. It is known that Ag is substituted for cations in BST, resulting in the production of holes.8, 17 Furthermore, it has been reported that Sb in Sb2Te3 is replaced by Ag, thereby generating holes.18 Trends in the lattice parameter have also been investigated for Ag doping in BST. In previous studies,8-9 the cell volume decreased at higher Ag concentrations due to the smaller size of Ag when compared to Bi or Sb. However, almost no variation in the cell volume was observed for the samples in this work (Table S1). Such a scenario could be ascribed to an Ag concentration that is sufficiently low so that a change in the lattice parameter is not detectable by XRD.


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Fig. 4 shows the Seebeck coefficient as a function of temperature. At room temperature, a decrease in the Seebeck coefficient was observed at higher Ag concentrations. For a highly degenerate semiconductor, the Seebeck coefficient can be expressed as follows:7 * =

+' (

$,)

-∗

'

$./

/$

(Eq. 2)

At room temperature, the observed decrease in the Seebeck coefficient at higher Ag concentrations is due to the increased carrier concentration after Ag doping. This finding is also supported by the electrical conductivity and Hall coefficient measurement results. A maximum in the value of the Seebeck coefficient (Smax) was reached for samples #1-#3, whereas the Seebeck coefficients for samples #4 and #5 monotonically increased over the entire range of testing temperatures (Fig. 4). Temperature-dependent Seebeck coefficient behavior with distinct Smax values has previously been observed in p-type Bi-Sb-Te compounds19-20 and is mainly due to the generation of minority carriers excited across the band gap.9 From the values of Smax and TS,max (the temperature of Smax), the band gap (Eg) can be evaluated from the expression,15, 21-22 12 = 2*

45 67, 45

(Eq. 3)

The value of Eg was estimated to be 0.19 eV for samples #1~#3; a similar value of 0.17 eV was obtained at almost the same composition in earlier work.15 Because Smax cannot be defined for samples #4 and #5, Eg could not be estimated. It is believed that Eg was not altered as a result of the ultra-low Ag concentration. The temperature at which Smax was obtained shifted to higher values as the Ag concentration increased (Fig. 4). In order to understand this behavior, the position of the Fermi energy was


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considered. Assuming single-band conduction, the Seebeck coefficient can be described according the following expression:15, 23 *=

( :; { } ,

(Eq. 4)

E

where Fn(η) is the Fermi-Dirac integral and . > = @C A . BC >DA , where f0 is the Fermi distribution, ξ is the reduced energy of the carriers, λ is the scattering parameter, which is assumed to be zero (acoustic phonon scattering) for all subsequent calculations, and η is the reduced Fermi energy (=Ef/kBT). Figure 5 shows the estimated Fermi energy at room temperature. While the value of the Fermi energy was evaluated under the isotropic single parabolic band assumption, findings related to the change in the Fermi energy with respect to Ag content can be effective for discussing the change in the Seebeck coefficient. The positive and larger value of Ef for samples #4 and #5 means that the Fermi energy is located far below the valence band maximum (VBM). As the amount of Ag is increased, Ef moves further down in the valence band. The values of Ef obtained for the samples range from -0.02 eV to 0.05 eV for samples #1 to #5, respectively. In other words, sample #1 is partially (moderate) degenerate, while sample #5 is fully (highly) degenerate (Fig. 5).23 The change in the temperature corresponding Smax will be discussed in relation to the change in the lattice thermal conductivity, where bipolar transport is considered. To investigate the change in the electronic structure, the effective mass was evaluated according to following equation (assuming that the Hall factor is unity):23-24 )

-∗ = (

./

[

G'

]⁄$

⁄

(Eq. 5)


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Using the reduced Fermi energy estimated from Eq. 4 and the measured carrier concentration, values of m* were calculated from Eq. 5; the results are shown in Fig. 6, along with m* values obtained from Eq. 2. It should be noted that the Eq. 2 was derived from a simplified model for highly degenerate systems, whereas Eq. 5 was formulated according to Fermi-Dirac statistics (FD statistics). When compared to the results calculated with FD statistics, the values of m* for samples #1~#3 were underestimated with degenerate statistics. In contrast, Eqs. 2 and 5 yielded similar m* values for samples #4 and #5. As mentioned above, the Fermi energies of samples #4 and #5 were well below the valence band maximum (VBM) and completely degenerate, which accounts for the consistency in m* from both expressions. The values of m* from FD statistics for samples #1~#3 are quite similar, about 1.35m0, which is larger than the 0.82m0 obtained in Ref. 9. The underestimation in the previous study may be attributed to the utilization of degenerate statistics. Thus, the results obtained with degenerate statistics can be misleading, as they show an increase in the effective mass for samples #2 and #3 when compared to that for sample #1. In other words, the use of Eq. 2 for highly degenerate semiconductors requires careful consideration. Interestingly, the relatively unchanged value of m* for samples #1~#3 means that the dispersion in the band edge is mostly unchanged for the specimens. Factors related to the electronic structure, namely Eg and m*, were mostly maintained for samples #1~#3. As the Fermi level moves down with larger Ag doping (samples #4 and #5), the effective mass exhibits a significant change due to the complicated electronic structure of BST.25 The power factor (σ·S2) is a combination function comprising the electrical conductivity and Seebeck coefficient. As the temperature was increased, the power factors of all samples tended to decrease, as shown in Fig. 7. While no large discrepancies in the power factor were observed from 25~50 °C, the values obtained at elevated temperature increased at higher Ag contents. The


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power factors of samples #4 and #5 were relatively constant or exhibited a slight decrease over the entire temperature range. Such behavior may be ascribed to the Seebeck coefficient, which exhibited no decrease at elevated temperature for these specimens. The thermal conductivity (κ) of the samples as a function of temperature is shown in Fig. 8. The room temperature κ values increased at higher Ag contents. Furthermore, the temperature dependence of κ varied depending on the sample. To verify the origin of the variation in κ, the electronic and lattice thermal conductivities were estimated by first calculating the Lorenz number with FD statistics as follows:23 (

J = ,

Method of Efficient Ag Doping for Fermi Level Tuning of

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