Spinodally Decomposed PbSe-PbTe Nanoparticles for High

Jul 9, 2016 - However, there exists no report on the n-type systems showing high figure of merits because of their intrinsically low Seebeck coefficie...
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Spinodally Decomposed PbSe-PbTe Nanoparticles for High-Performance Thermoelectrics: Enhanced Phonon Scattering and Unusual Transport Behavior Min-Seok Kim,†,‡ Woo-Jin Lee,‡ Ki-Hyun Cho,‡ Jae-Pyoung Ahn,§ and Yun-Mo Sung*,‡ †

Computational Science Center, Korea Institute of Science and Technology 5, Hwarang-ro 14-gil, Seongbuk-gu, Seoul 136-791, Korea Department of Materials Science and Engineering, Korea University, Anam-Dong 5-1, Seoul, 136-713, Korea § Advanced Analysis Center, Korea Institute of Science and Technology 5, Hwarang-ro 14-gil, Seongbuk-gu, Seoul 136-791, Korea ‡

S Supporting Information *

ABSTRACT: Dramatic enhancements in the figure of merit have been obtained in bulk thermoelectric materials by doping, band engineering, and nanostructuring. Especially, in p-type thermoelectrics, high figure of merits near 2.0 have been reported in a few papers through the reduction in lattice thermal conductivity and the advancement in power factors. However, there exists no report on the n-type systems showing high figure of merits because of their intrinsically low Seebeck coefficients. Here, we demonstrate that a nanostructured bulk n-type thermoelectric material that was assembled by sintering spinodally decomposed lead chalcogenide nanoparticles having a composition of PbSe0.5Te0.5 reaches a high figure of merit of 1.85. The spinodally decomposed nanoparticles permit our thermoelectric material to have extremely low lattice thermal conductivity and a high power factor as a result of nanostructuring, electronic optimization, insertion of an impurity phase and phase change in local areas. We propose that this interesting concept would be one of the promising approaches that overcome limitation arising from the fact that most parameters in the figure of merit are closely correlated. KEYWORDS: thermoelectrics, spinodal decomposition, nanoparticles, spark plasma sintering

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Except for superlattice structures, only a few nanostructured bulk materials with maximum ZT values of 2.0 or higher have recently been reported. Biswas et al.20 reported a ZT value of 2.2 via all-scale hierarchical architectures in p-type PbTe with 4 mol % SrTe. They successfully reduced κL by scattering broadwavelength phonons at all scales, i.e., from the atomic scale to the mesoscale. This was achieved by fabricating architectures with atomic-scale dopants, heterogeneous nanoscale precipitates, and mesoscale grain boundaries, all of which resulted in maximum phonon scattering, i.e., a low κL value. Gelbstein et al.47 obtained a ZT value of approximately 2.2 using phase separation and transition phenomena of Ge0.87Pb0.13Te. Moreover, they stabilized thermoelectric properties by subsequent heat treatments and these sophisticated metallurgical controls

t is estimated that approximately 80% of the global energy consumption is based on fossil fuels, which have limited reserves, and this high energy dependence has led to the development of renewable-energy technologies.1 Of all the renewable-energy technologies, thermoelectric (TE) materials, which can convert waste heat into useful electrical power, could play a significant role in reducing the dependence on fossil fuels because of some approaches that have been developed, such as the discovery of new TE materials,2−7 the use of functionally graded materials,53 nanostructuring,8−10 band engineering,11,12,14 resonant doping,12−14 and energy filtering.6,15−18 As a result of these techniques, the dimensionless figure of merit, ZT = S2σT/κT,19 which represents the efficiency of a TE material, has been improved considerably over the last two decades. Here, S, σ, T, and κT represent the Seebeck coefficient, electrical conductivity, absolute temperature, and total thermal conductivity, which is the sum of the electronic and lattice thermal conductivities (κE and κL, respectively). © 2016 American Chemical Society

Received: June 5, 2016 Accepted: July 9, 2016 Published: July 9, 2016 7197

DOI: 10.1021/acsnano.6b03696 ACS Nano 2016, 10, 7197−7207

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Generally, in lead chalcogenide TE materials, μ tends to decrease with increasing temperature and then it gradually approaches a certain value at the S-saturated temperature. By using spinodal decomposition, Androulakis et al.30 reported remarkably reduced κL values for the PbTe/PbS system. Similarly, we used spinodal decomposition to reduce the κL value. However, the resulting behaviors of μ in the two studies are quite different. This difference might originate from the different starting materials used. In the PbTe/PbS system,30 Pb, Te, and S elements were used as starting materials and the spinodal decomposition was induced during the fabrication of the PbTe/PbS thermoelectric bulks. In our case, the spinodally decomposed PbSe/PbTe NPs were used as starting materials and the thermoelectric bulks were obtained using these NPs by the sintering. It seems to be interesting that the spinodal decomposition induced during the heat treatment30 and the use of the spinodally decomposed NPs as starting materials derive the independent results.

based on a phase diagram lead to a stable ZT value of approximately 2.0. In addition, through band-structure engineering and the suppression of the bipolar effect, Zhao et al.21 obtained a ZT value of approximately 2.0 with a Na-doped PbTe system containing 6 mol % of MgTe. Wu et al.22 also recorded a ZT value greater than 2.0 over a very wide temperature range (673−923 K). They explained that this phenomenon was the result of tuning the carrier concentration in addition to the band-structure engineering and hierarchical structure. Wang et al.23,24 also obtained a ZT value of 2.0 with Na-doped PbTe samples. Just recently, Wu et al.46 recorded a ZT value of 2.3 using “abnormal saturation” phenomena. To the best of our knowledge, the aforementioned studies, which focus on PbTe-based material, are the only studies that have reported isotropic ZT values of around 2.0 or more.25 However, there have been no studies on any n-type TE systems that have similarly high ZT values despite the structure of TE devices, which requires n- and p-type TE elements to be connected electrically in series and thermally in parallel.26 Hsu et al.3 have reported a ZT value of 2.2 in an n-type semiconductor of AgPbmSbTe2+m; however, a lower ZT value was confirmed in a subsequent report54 because of macroscopic thermoelectric inhomogeneities and nowadays the ZT values of 1.5−1.7 are recognized in those kinds of materials.55 One of the main reasons for this is that n-type lead chalcogenides have low intrinsic S values compared to p-type materials because of the relatively low number of degenerate valleys.11,27 In addition, there have not been any successful cases of resonant doping or distortion of the density of states (DOS) in n-type TE materials, whereas Heremans et al.13 reported a doubling of ZT via the introduction of Tl impurity levels inside the valence band of PbTe. In the present study, we found that the molar ratio of PbSe to PbTe is important for achieving a high ZT value. By using spark-plasma sintering (SPS) with the prepared powders, we were able to produce nanostructured bulk specimens. The SPStreated PbSe/PbTe samples exhibited markedly reduced κT values and increased S values compared to that of the SPStreated PbSe samples. Moreover, we discovered an unusual transport behavior in the SPS-treated PbSe/PbTe samples. As a result, we recorded a ZT value of approximately 1.85 at 623 K, which is the highest reported value for n-type TE structures to date. We believe that this breakthrough is the result of applying spinodally decomposed NPs to TE materials based on the two following reasons (note that the world of “decomposed” does not mean that PbSe/PbTe NPs convert their phase from the solid-solution to the spinodal during the synthesis. Here, we want to emphasize that the synthesized NPs have the spinodally decomposed regions). In contrast to conventional powderprocessing techniques, the bottom-up assembly of NPs allows TE materials to have grain sizes that are comparable to the predicted optimum conditions for hindering phonons, especially those with long mean free paths. According to firstprinciples calculations,18,36 widespread inhomogeneities that are smaller than 1 μm might be more effective at reducing κL. However, the aforementioned studies reporting maximum ZT values of 2.0 or greater do not satisfy this condition.20−24,46,47 In our SPS-treated samples, the grain sizes ranged from 100 to 700 nm, which corresponds to the distinguishable conditions from the aforementioned studies for scattering phonons with long mean free paths. On the other hand, our SPS-treated PbSe/PbTe samples exhibited an unusual transport behavior, i.e., the mobility (μ) began to increase at a certain temperature.

RESULTS AND DISCUSSION Synthesis of Spinodally Decomposed NPs. Figure 1a shows the X-ray diffraction (XRD) patterns of the as-prepared PbSe/PbTe NPs and SPS-treated bulk material. The XRD peaks of the PbSe/PbTe NPs indicate the presence of two independent phases, PbSe and PbTe, and both sets of peaks are shifted to lower 2θ angles by approximately 0.3−0.4°. These shifts in the peaks appear to be caused by lattice expansion because of the excess Pb atoms.28 A quantitative analysis of the crystallinity of the phases in the NPs was performed using a multipeak separation program (Rietveld tool included in JADE 2010; Materials Data Incorporation), and the molar ratio of PbSe to PbTe is estimated to be 55 to 45. Usually, the two separated peaks indicate that synthesized NPs are composed of individual PbSe and PbTe NPs or PbSe/PbTe core/shell NPs because spinodal decomposition is generally difficult to induce in NPs, which is attributed to the reduction of the spinodal gap.29,48 However, by using transmission electron microscopy (TEM; Figure 1b, we observed NPs with a size of 40−60 nm that manifested continuous bands. Both PbSe and PbTe can initiate their respective nucleation and growth, spinodal decomposition, or solid-solution formation depending on the temperature and composition (Figure S1a, Supporting Information). The main feature of spinodal decomposition is that each phase has continuous clusters of its most extreme composition and exhibits nonspherical bands with high connectivity. The continuous bands in the PbSe/PbTe NPs are evidence of phase-separated nanocrystals, and this is supported by the high-resolution TEM image (Figure 1c). Moreover, a fast Fourier transform (FFT) image of Figure 1c and its inverse illustrate that local compositional fluctuations exist inside a single PbSe/PbTe NP (Figure 1d and e, respectively). For clarity, the inverse images extracted from the other spots in Figure 1d are depicted in the Supporting Information (Figure S1b). In this paper, we do not focus on the spinodal decomposition phenomenon itself as it is well described in other papers.30,31 In addition, we are currently preparing an in-depth paper that is related to spinodal decomposition in single NPs, and we have already obtained interesting results using in situ high-temperature XRD. More high-magnification TEM images are given in the Supporting Information (Figure S2) and the defects in a single NP are identified. 7198

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Figure 1. Spinodally decomposed PbSe/PbTe NPs. (a) XRD patterns of the as-prepared PbSe/PbTe NPs and SPS-treated bulk material. The reference XRD patterns (JCPDS 78−1903 and 78− 1905) for cubic PbSe and PbTe phases are presented at the bottom as red and blue vertical lines, respectively, and the first major peak of each phase is indicated at the top with dashed orange lines. A.U. = arbitrary units. In the XRD pattern of the SPS-treated material, the “*” denotes diffractions from a Pb phase. The explanations for the “shoulder” and “∼0.3−0.4°” comments are in the text. (b) Lowmagnification TEM image of the spinodally decomposed PbSe/ PbTe NPs. (c) High-magnification TEM image of a single spinodally decomposed PbSe/PbTe NP. The defects inside a single NP are given in the Supporting Information (Figure S2). The scale bars in (b) and (c) are reconstructed for clarity. (d,e) FFT image of (c) and its inverse image (e) are illustrated. For clarity, the inverse image was extracted from one spot. Inverse images of the other spots are depicted more clearly in the Supporting Information (Figure S1b).

Figure 2. Structural analysis of the bulk nanostructured PbSe/ PbTe. (a) Low-magnification TEM image showing various grains that are 100−700 nm in size. Two distinct regions are indicated with “A” and “B”. (b) Magnified image of grain “B” in (a). It contains various phases, such as α, β, and γ phases, precipitates with Moiré fringes (white dashed circles), and a background host phase. (c−f) High-magnification TEM images of the various phases in (b): (c) α, (d) β, (e) precipitates, and (f) γ. All scale bars were reconstructed for clarity. More evidence of the imperfections in the crystallinity because of an excess of Pb atoms (Figure 2c) is shown in the Supporting Information (Figure S3).

TEM Study of the Nanostructured Bulk PbSe/PbTe. To reduce κL efficiently, a uniform distribution of scattering centers is critical for preventing heat-carrying phonons, as well as various scattering centers over a wide range of length scales. Kanatzidis et al. have reported the most successful approaches for the reduction of κL. They produced various scattering sources to block broad-wavelength phonons at all scales.20,22 Interparticle spacing is also an important factor to reduce κL. For example, the large interparticle spacing or nonuniform distribution of scattering sources did not exhibit observable κL reduction.12 Using spinodally decomposed NPs, we can successfully create a microstructure that suppresses κL more than in any previous studies. Figure 2a shows that the SPStreated bulk material is composed of grains that are approximately 100−700 nm in size, which may be the result of grain growth during the sintering process, and they are tightly bonded to each other. The grains reveal two distinct sections, which are marked as “A” and “B”. The region of “A”

can convert its morphological features from textured to smooth and the phase variation during the heating and cooling processes is reversible as shown in Figure 4c. The region of “B” does not convert its morphological features depending on temperature and it is stable until 623 K. The details of “A” will be covered later alongside the electronic transport behavior. Through multiple TEM images taken at different spots, we estimate that the volume fraction of “A” is approximately 25%. The arbitrarily chosen grain “B” is magnified in Figure 2b and it consists of various phases, including α, β, and γ phases, precipitates with Moiré fringes (these occur due to crystal overlap between a background host phase of PbSe0.5T0.5 and precipitates of PbSe1−xTex (0.8 < x < 1.0)), and a background host phase. From the XRD pattern (Figure 1a), we confirmed that the SPS-treated bulk material is composed of a solid solution of PbSe and PbTe, and the composition is approximately 50:50 (i.e., PbSe0.5Te0.5). There also exist secondary peaks of a Pb phase, and these compositional 7199

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Figure 3. Thermoelectric transport properties. (a) Temperature-dependent total thermal conductivity (κT) and lattice thermal conductivity (κL) for the nanostructured bulk PbSe and PbSe/PbTe materials. The κL values were estimated by subtracting the electronic thermal conductivity (κE) from the measured κT values at each temperature. The κE values were calculated with the equation of κE = LσT (where L is the Lorenz number, σ is the electrical conductivity, and T is the absolute temperature, Supporting Information). The purple arrow in (a) indicates the temperature at which the κT values of the PbSe/PbTe samples begin to increase. The plot includes the κT values for one cycle of measurements for the PbSe samples and two cycles for the PbSe/PbTe samples. (b) Negative Seebeck coefficient (S) as a function of the temperature. The negative values indicate that both TE materials have n-type characteristics. The green dashed lines serve as a guide for the eye. (c) Electrical conductivity (σ) as a function of the temperature. The purple arrow in (c) indicates the temperature at which the σ values of PbSe/PbTe samples begin to increase. (d) Power factor (S2σ) as a function of the temperature.

well with the XRD data in Figure 1a because the XRD peaks in the SPS-treated bulk material have shoulders on the left-hand side that are closer to PbTe than PbSe. The variable contrast within the precipitates is caused by the Moiré patterns that are produced by the lattice mismatches between the superimposed lattices of the host and precipitates.33 All of the aforementioned precipitates appear to be generated from nucleation and growth phenomena. However, there are specific regions that are approximately 40−60 nm wide and 100−150 nm long, which consist of almost parallel stripes, as shown in Figure 2f. The stripes are spaced by approximately 2−3 nm and are a common characteristic of spinodal decomposition.30 Interestingly, our samples possess various scattering sites that vary from the subnm to 700 nm in size, and this is a major factor in the markedly reduced κL (this is covered in detail later in the paper). Thermoelectric Transport Properties. The first measurement of the TE properties of the SPS-treated bulk PbSe/PbTe resulted in an unexpectedly high ZT value above 2.2 at 823 K. However, the second measurement of the first-tested sample resulted in a lower ZT value of 1.4 because of the increase in κL and decrease in the power factor (S2σ). And we found that the TE properties deteriorated as the measurements were repeated.

characteristics are consistent with the additional chemical analyses performed with inductively coupled plasma-atomic emission spectrometry (ICP-AES; Pb:Se:Te = 54:21.2:24.8 (mole fraction), there could exist an error of ±5%.). From the XRD patterns and results obtained with the energy-dispersive X-ray spectrometer (EDS) attached to the FE-TEM setup, we concluded that the host phase is a solid solution of PbSe1−xTex (0.45 < x < 0.5). Inside each grain, there is a region of an α phase with a darker contrast than that of the host phase, as shown in Figure 2c. These regions have a slightly higher Pb contents and exhibit imperfections in the crystallinity or line defects (Figure S3, Supporting Information), which is in good agreement with results in previous papers.20,32 In the β regions, we were able to identify precipitates as small as 2−10 nm with high-magnification TEM, as shown in Figure 2d, which are indiscernible in Figure 2b. A uniform distribution of larger precipitates also resides in the grains, as indicated by the dashed circles in Figure 2b. A magnified image (Figure 2e) shows that these nanostructures are endotaxially embedded in the host phase and are approximately 10−30 nm in size. EDS analysis revealed that these precipitates have a higher Te content than the surrounding host material, and their chemical composition is estimated to be PbSe1−xTex (0.8 < x < 1.0). This result agrees 7200

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ACS Nano Once the temperature rose above ∼650 K, which corresponds to the temperature of the immiscibility boundary between the spinodal and the solid-solution regions (see Figure S1a of the Supporting Information), the samples did not hold any stability and we could not recognize any tendency between each measurement. Therefore, the TE properties were only measured up to 623 K and the samples exhibited the stability in this temperature range. For a high ZT value, some requirements must be met (note that these requirements result from our experimental trials and errors). One such requirement is that the spinodally decomposed NPs should have a composition ranging from 45:55 to 50:50 (PbSe and PbTe in this case). This can be easily calculated from the XRD patterns by comparing the experimental patterns to references or using an analytical software package like Jade (Jade Software Cooperation), as mentioned above.34 Through experimental trials and errors, we found that the different composition range did not exhibit the reversible phase transition as shown in Figure 4c, and therefore, the specific increase of μ also did not occur. The other requirement is that the SPS-treated samples should not have an excess of atoms, except for Pb atoms. Inappropriate sintering conditions can occasionally cause the formation of a Se phase in the samples. Whenever we detected Se diffraction peaks in the XRD patterns, σ was reduced because of the bipolar effect.6,35 When the aforementioned conditions were met, a stable ZT value over 323−623 K was achieved. Repeating the TE measurements over the aforementioned temperature range guarantees the consistency of the TE parameters, and one data cycle for the PbSe NPs and two data cycles for the bulk PbSe/PbTe are depicted in Figure 3. The data presented for the PbSe/PbTe was acquired via a round robin test to prevent any possible inaccuracies that may arise from the characteristics of our measuring devices. (For each cycle, the measurements were performed more than 10 times per one sample. Various samples were tested and average values are presented in the result. The first cycle is performed in our laboratory and the other is performed in Yonsei University, Seoul, Korea.) The TE properties were measured at around 323, 373, 423, 473, 523, 573, and 623 K while using regular intervals of 5, 10, and 15 K per cycle, with Figure 3 showing the average values. The sample preparation for each cycle of measurements is schematically shown in the Supporting Information (Figure S4). The same crystallographic orientations were chosen for the measurement of each TE parameter. Recently, the extremely high isotropic value of ZT (ZTmax) of approximately 2.2 has been reported for a p-type material at 915 K. The main reason for this high ZT value is the remarkably low κL, which was achieved with a panoscopic approach.20,22 In the present study, we reduced κL beyond what was achieved in the aforementioned studies, even at the relatively low temperature of 623 K, and reaffirmed the effectiveness of the panoscopic approach. Two possible reasons were suggested for the origins of such low κL values: (1) the grains are smaller and (2) phonon scattering occurs at all scales. In the two aforementioned papers,20,22 the average size of the grains was close to 1 μm. In other words, grains that were larger than 1 μm were present in the samples (both PbTe and PbS were considered by Wu et al.22). However, the SPS-treated PbSe/PbTe samples in the present study had an average grain size of 400 nm and grain-size distribution of 100−700 nm, which is the result of using spinodally decomposed NPs prepared with a bottom-up chemical synthesis. Naturally, it is hard to obtain TE specimens with grains of less than 1 μm

because the mechanical milling of TE ingots can only produce powder with a particle-size distribution of several hundreds of nanometers (especially in the lead chalcogenide system). According to the calculations of κL for PbSe, PbTe, and their alloys, most of the mean free path (MFP) values remain within 1 μm.18,36 Biswas et al.20 also pointed out that 20% of the κL value of PbTe is contributed by phonon modes with MFPs of 0.1−1 μm. The grain-size distribution of the SPS-treated bulk structures is similar to these values, as shown in Figure 2a. With respect to reason (2), low ZT values are obtained if there is a dense and uniform distribution of phonon scattering centers over a wide range of length scales. Figure 2 shows that the SPStreated PbSe/PbTe samples carry various phonon scattering channels at a variety of scales, such as the point and line defects at the atomic scale (Figure 2c and Supporting Information Figure S3); spinodal patches and small precipitates that are 2−3 nm and 2−10 nm in size (Figure 2d and f); large precipitates that are 10−30 nm in size (Figure 2b and e); spinodal bands (Figure 2b and f) with a width and length of 40−60 nm and 100−150 nm, respectively; and grains that are 100−700 nm in size (Figure 2a). It is believed that other decisive parameters that determine the effect of phonon scattering are the interparticle spacing between the precipitates and morphology of the precipitates.37,38 In the case of this study, the interparticle distance is approximately 5−15 nm for the small precipitates and 10−30 nm for the large precipitates, as shown in Figure 2b and 2d. Moreover, these precipitates are widely spread throughout the host phase and have a high population density and three-dimensional structure. All of these factors result in PbSe/PbTe-based TE materials with very low κT values compared to that of pure PbSe. Figure 3a shows the κT values of the SPS-treated bulk materials synthesized with PbSe and PbSe/PbTe NPs. A κT value of 1.91 W/mK was obtained with the PbSe samples at 323 K, and it decreases to 1.22 W/mK at 623 K. In contrast, a value of 1.18 W/mK was obtained with the bulk PbSe/PbTe at 323 K, and it decreases to 0.84 W/mK at 573 K and then increases to 0.88 W/mK at 623 K in both measurement cycles. This increase at higher temperatures will be discussed later. The κL values were estimated by subtracting κE from κT (see the Supporting Information for further details). Figure 3a shows that the κL values for PbSe decrease from 1.34 W/mK at 323 K to 0.58 W/mK at 623 K, and those for the bulk PbSe/PbTe are smaller than expected. The values decrease as the temperature increases and reach 0.40 W/mK at 623 K. This κL value is so low that it even reaches the amorphous limit of ∼0.32 W/mK (at high temperature) for the bulk PbTe system.45 Here, another factor to notice is that a specific heat capacity (Cp) value of the bulk PbSe/PbTe is slightly higher than a theoretical value according to Dulong−Petit rule (Supporting Information Figure S8). Materials with nanograins frequently exhibit higher Cp values than estimated ones because they have large surface area.33 And, therefore, the surface contribution to a Cp value needs to be considered because an underestimated Cp value will immediately affect a ZT value (κ is directly connected with Cp by the equation of κ = DCpd). Despite correction of Cp, the PbSe/PbTe samples still reveal the exceptionally low κL value. This may be one of the lowest κL value reported to date in a bulk lead chalcogenide system, especially at the relatively low temperature of 623 K.20−22,30 (possible effect of porosity is also accounted for in the thermal conductivity, see Supporting Information). Figure 3b shows that the S values of all of the NP-based composites have negative values, i.e., they are n-type TE 7201

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ACS Nano 2 π 2 kB ⎡ 1 dn(E) S 1 dμ(E) ⎤ = + ⎢ ⎥ μ(E) dE ⎦ E = E T 3 q ⎣ n(E) dE

materials. This n-type conduction is due to the deficiency of Se and Te atoms (a more detailed explanation is provided in the Supporting Information). The PbSe specimens exhibit a linear increase in S (85 to 166 μV/K) until a temperature of 473 K is reached, and then the rate of increase of the S values decreases above 473 K. Finally, S reaches a value of 188 μV/K at 623 K. In contrast, the S values of the PbSe/PbTe samples exhibits two linear regions that are divided at 523 K, i.e., S increases steeply from 103 to 225 μV/K as the temperature increases from 323 to 523 K, respectively, and then S gradually increases to 238 μV/K at 623 K. The Hall carrier densities for each sample were measured to estimate the temperature variance of S. The carrier densities of the PbSe and PbSe/PbTe samples were found to be 1.9 × 1019 and 1.5 × 1019 cm−3 (Supporting Information Figure S6), respectively, and these values are relatively independent of temperature. Further, we measured high-temperature hall-effect to confirm the thermal stability after annealing at 623 K for 24 h. And we concluded that there are no discernible mobility changes in both PbSe and PbSe/PbTe samples (Supporting Information Figure S6). Using the Kane band model, the S values of each sample were estimated and the values were compared to the experimental data (see Figure S5 of the Supporting Information). The specific calculation method was described in detail by Wang et al.,45 and it is summarized in the Supporting Information. In the case of PbSe, the experimental S values are initially similar to the estimated values. The experimental values increasingly deviate as the temperature increases to 473 K, and then approach the estimated values at 623 K because of the saturation of S, which begins at 473 K. Unlike the PbSe results, the bulk PbSe/PbTe shows highly enhanced S values compared to the estimated values. Although the rate of increase of S decreases above 523 K, the difference between the S values is still considerable at 623 K, and this is further evidence of the high ZT values of our PbSe/PbTe samples. The enhanced S values of PbSe/PbTe samples compared to that of PbSe samples appear to arise from the energy barriers at the interfaces, including the precipitate/ matrix interfaces and grain boundaries. Moreover, some reports showed that impurity phases (Pb impurity phases in our case) can influence the transport properties within a semiconductor.15,44,39,49 An in-depth discussion of this is included in the next section. It should be emphasized that the increase in σ for the bulk PbSe/PbTe samples starts at 523 K, which is identical to the temperature at which the rate of increase of S decreases, as shown in Figure 3c (this is discussed in detail later). The power factor (PF = S2σ) values for both systems as a function of the temperature are shown in Figure 3d. Despite the higher S values of the bulk PbSe/PbTe samples, its PF values are smaller than that of the PbSe samples until 573 K is reached, which is because of the lower σ values (Figure 3c). However, the values are reversed near 600 K because of the unusual variation in σ. The PF values of the PbSe and PbSe/ PbTe samples increase from 6.6 and 4 μW/cm·K2 at 323 K, to 23.6 and 27.5 μW/cm·K2 at 623 K, respectively. Unusual Behavior of the Electron Mobility. The nanostructured bulk PbSe/PbTe materials exhibit two distinct regions with different gradients in the plot of S as a function of the temperature (T). Regarding the ratio of S/T, there is a useful equation that originates from the Bethe−Sommerfeld expansion of the Mott formula:13,14

F

where q, kB, n(E), and μ(E) are the charge, Boltzmann constant, energy-dependent carrier concentration, and mobility at the Fermi energy (EF), respectively. Therefore, the S/T values can vary according to two terms: [1/n(E)][dn(E)/dE] and [1/μ(E)][dμ(E)/dE]. According to previous studies, it is difficult to control the former term because of the distortion of the electronic DOS or introduction of resonant levels inside the conduction band of the nanostructured PbSe/PbTe materials, which is unlikely.6,12−14 It is probable that the deviation of S/T at around 523 K is due to the variation of [(1/μ(E)][(dμ(E)/ dE)]. The Hall measurements we performed confirmed our hypothesis. The Hall carrier densities according to temperature are nearly identical in both samples, and therefore, the [1/ n(E)][dn(E)/dE] term cannot cause the enhancement of S. An important factor that affects the [1/μ(E)][dμ(E)/dE] term is the scattering parameter (r) because r can change the relaxation time (τ) and μ according to the following equations (dominant mechanism of scattering can be inferred from r. Generally, TE materials exhibit r ≈ 0 because the electron scattering by phonons is dominant at high temperature, and therefore, μ tends to decrease as temperature increase. However, μ of our PbSe/PbTe samples exhibit the increase with increasing temperature from 523 K. More detail explanations for r are well presented in following references): τ = τo·Er−1/2 and μ = qτ/m, respectively (where E is the energy and m is the effective mass of the carrier).12,14,15,39,44 In particular, Heremans et al. accomplished thermopower enhancements in both PbTe with Pb precipitates44 and nanostructured PbTe.15 The factors that induce these enhancements are the energy barriers provided by precipitates and grain boundaries (according to their discussions, these energy barriers can change the scattering parameter (r) and it leads to the enhanced S value). These systems are similar to those in the present study because PbSe/ PbTe samples have an excess of Pb atoms, Te-rich PbSe/PbTe precipitates with a wide size distribution, and many grain boundaries. Using effective medium theory, Day et al.49 successfully separated conductivity and mobility of independent phases and showed an impurity phase can severely influence TE transport properties. About an application of the effective medium theory to multiple-component TE materials, Gelbstein et al.50,51 explained well in some papers. The energy dependence of μ also affects σ and κE (σ = nqμ). The increase in σ with increasing temperature, which is mainly caused by an alteration of scattering modes, is unusual and very few studies have reported this phenomenon with nanostructured TE materials. This unusual behavior of μ has been found in PbTe that was conanostructured with both Pb and Sb precipitates.39 Together with the above reports,15,44,49 this article suggests that Pb precipitates with or without additional atoms can play a significant role in the Seebeck coefficient and the electrical transport. From the XRD pattern, as shown in Figure 1a, we also identified the precipitates of Pb and they would also play a significant role in our system. However, at this stage, it is hard to recognize which precipitates lead to a high power factor (S2σ). Our system has a lot of interfaces (energy barriers) like the excess of Pb atoms, Te-rich PbSe/PbTe precipitates with a wide size distribution, and many grain boundaries. We guess that all the interfaces simultaneously contribute to the high 7202

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Figure 4. Unusual behavior of the electron mobility. (a) Temperature-dependent electrical conductivity (σ, left axis) and electronic thermal conductivity (κE, right axis) for the bulk nanostructured PbSe/PbTe. Both plots exhibit the similar trend. (b) In situ high-temperature XRD data. Only a small sample of the data is plotted and no distinct phase changes can be seen during the heating and cooling processes. The purple arrows in (b) indicate the heating and cooling processes. (c) In situ high-temperature TEM images. Only a few are images are shown that demonstrate the observable phase variation that occurs in certain local grains (∼25% in volume, the “A” grains in Figure 2a). The phase variation during the heating and cooling processes (green arrows) is reversible, as mentioned in the text. All scale bars were redrawn for clarity and indicate a size of 200 nm.

S values is still considerable at 623 K. The main difference between two TE materials is the presence of Pb precipitates. As shown in Figure 1a, the PbSe/PbTe TE materials have extra Pb precipitates. However, XRD data from the pure PbSe do not show such precipitates as shown in Supporting Information Figure S7. On the basis of our experimental results, it seems that Pb elements are more soluble in PbSe than PbSe/PbTe (note that Se has higher electronegativity and the carrier concentration is higher in PbSe than Pbse/PbTe). Pb atoms, which could not dissolve in PbSe/PbTe, would be the key factor to increase the S value as shown in mentioned references.15,44,39,49 More recently, σ values that were 10-fold higher than that of pure PbTe and PbS NPs have been achieved in bulk samples that were synthesized with PbTe-PbS core− shell NPs.33 A modulation-doping approach, which requires at least two different types of nanograins to exist, is considered an effective way to improve the power factor by mainly increasing μ.40,41 Regarding our PbSe/PbTe samples, σ exhibits a steep increase at around 523 K. This result is supported by the plot of κT (Figure 3a), which shows an increase at 573 K while the κL values steadily decrease as the temperature increases. In Figure 4a, the σ and κE values of the PbSe/PbTe nanocomposites are plotted as a function of the temperature and similar trends are exhibited, i.e., both parameters rapidly increase from 523 K. Using in situ high-temperature XRD and TEM, we attempted to trace the factor governing the escalation of μ, as shown in Figure 4b and c. The in situ XRD measurements allowed the monitoring of any noticeable phase transition when the sample

power factor because there is no observable transport behavior in the pure PbSe. Although it is difficult to recognize what the most critical factor among the various energy barriers, we try to explain the reason for the enhanced S value in PbSe/Te compared to PbSe. The carrier densities of the PbSe and PbSe/ PbTe samples were found to be 1.9 × 1019 and 1.5 × 1019 cm−3 (Supporting Information Figure S6), respectively. And it is normal the lower carrier concentration lead to the higher S value.26,50,51 Especially, Wu et al.52 reduced the carrier concentration in Ge0.87Pb0.13Te upon the introduction of 3 mol % Bi2Te3 and resulted in the enhanced S value. They verified that the solubility of PbTe in GeTe can be altered according to the presence of Bi2Te3. The estimated S values of each sample also exhibit this tendency as shown in the Figure S5 of the Supporting Information. However, it seems that there are other effects of increasing the S values in our system except for the optimization of the carrier concentration. We compared the estimated S values with the experimental data in PbSe. We performed the same thing for PbSe/PbTe. However, the derived results are quite different despite application of the same Kane band model. In the case of PbSe, the experimental S values are initially similar to the estimated values. The experimental values deviate as the temperature increases to 473 K, and then approach the estimated values at 623 K because of the saturation of S, which begins at 473 K. Unlike the PbSe result, the bulk PbSe/PbTe shows highly enhanced S values compared to the estimated values. Although the rate of increase of S decreases above 523 K, the difference between the 7203

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ACS Nano was heated and cooled between 323 and 623 K in intervals of 50 K. However, no distinct phase transition was observed. The XRD patterns at 573 and 623 K during the heating phase and that at 573 K during the cooling phase are shown in Figure 4b. On the other hand, instead of a bulk-scale phase transition, there could be discernible phase variations in individual grains. Through the TEM analysis of the nanostructured bulk PbSe/ PbTe, we discovered grains (approximately 25% in volume, denoted “A” in Figure 2a) that exhibit variable levels of contrast as the temperature changes, and these grains appear to undergo spinodal decomposition. One of these grains was carefully investigated with in situ TEM, as shown in Figure 4c. When the temperature is raised from 550 to 600 K, the spinodally decomposed grain with a nonuniform texture and apparent boundary loses its morphological features because of the transition to a solid-solution region (refer to Figure S1a of the Supporting Information), but the lost features are recovered when the grain is cooled from 600 to 550 K. This interesting phenomenon could correlate with the increase in both σ and κE above 523 K. In other words, μ increases at around 523 K because of reduced number of scattering sources, which is the same temperature where κT increases and κL gradually decreases (Figure 3a). The instability in the TE properties above 623 K appears to originate from the absence of the aforementioned reversibility. The diffusion of elements occurs from the “A” grain to the “B” grain, with the “A” grain eventually losing its grain boundary above 623 K, which corresponds to the temperature for the phase transition from the spinodal to solid-solution structures, as depicted in Figure S1a of the Supporting Information. Up to 623 K, the grain boundary of the “A” region is maintained, even though there are local fluctuations in the phases, as shown in Figure 4c. Using the spinodally decomposed PbSe/PbTe NPs, the highest ZT value obtained was approximately 1.85 at the relatively low temperature of 623 K for n-type systems. Figure 5d shows that the ZTmax of the PbSe/PbTe material exhibits an improvement of approximately 154% compared to that of the nanostructured bulk PbSe materials (ZTmax = 1.2 at 623 K). The major factors that produce the high ZT values in the nanostructured bulk PbSe/PbTe are depicted in Figure 5a−c. First, we developed an effective way to block the short-, medium-, and long-wavelength heat-carrying phonons (Figure 5a), which in turn resulted in a very large reduction in κL. Scattering centers, such as point and line defects, small and large precipitates, spinodal bands, and grains, exist with an optimal distribution when considering the size, density, and interparticle spacing of such scattering centers. Second, the Seebeck coefficient is enhanced via the variation of the scattering mode (r). Energy barriers among host phase, precipitates, and grain boundaries permit this effect, as illustrated in Figure 5b. Finally, the sintering of the spinodally decomposed PbSe/PbTe NPs results in an unusual electronic transport behavior because specific local grains can reversibly transform between spinodal and solid-solution structures as shown in Figure 5c. This distinct structure would be hard to obtain with conventional powder processing and a subsequent sintering process because there is the limitation of forming mesoscale grains (around 1 μm in size) instead of the omnipresent phonon scattering centers over a wide range of length. The existence of two different phases in a single NP allows results in not only effective scattering centers but also separated “A” and “B” grains, which cause σ and S to increase with temperature.

Figure 5. Mechanisms that contribute to the high ZT value (approximately 2.0) of the n-type system at the relatively low temperature of 623 K. (a) Schematic diagram illustrating the alllength-scale hierarchy in the TE material. By using spinodally decomposed PbSe/PbTe NPs, it is possible to create mesoscale grains that are smaller than 1 μm and omnipresent phonon scattering centers over a wide range of length scales (sub-nm: solidsolution and excess Pb atoms; 1−100 nm: small and large precipitates, and spinodal bands and their patches; 100−700 nm: grain boundaries). The schematic figure was redrawn with the aid of Figure 4a in ref 27. (b) The various scattering sites. Pure PbTe and PbSe0.5Te0.5 were selected for the precipitates and host phase, respectively, and the band-edge difference (CBE = conductionband edge) between them can be calculated according to ref 42. A number of energy barriers in the host phase, precipitates, and grain boundaries increase the Seebeck coefficient (S) through changing the scattering mode (r). (c) Unusual electronic transport behavior. Specific local grains can reversibly transition between spinodal and solid-solution structures, which increase the mobility (μ). (d) ZT as a function of the temperature for the bulk nanostructured PbSe and PbSe/PbTe materials. Owing to the mechanisms described in (a), (b), and (c), the highest ZT value of approximately 1.85 is achieved with the n-type PbSe/PbTe TE material at the relatively low temperature of 623 K. The standard deviation in the data because of variations in the samples and measurements is indicated by the error bars (excluding the errors specified in the equipment manuals).

CONCLUSION In this study, we discovered phenomena that result in a high ZT value when using spinodally decomposed PbSe/PbTe NPs. We obtained a ZT value of approximately 1.85 at the relatively low temperature of 623 K, and to the best of our knowledge, this is the highest recorded value for an n-type system. The achievement of such high ZT values can only stem from the bottom-up assembly of spinodally decomposed NPs. We anticipate that these phenomena will provide possible approaches that overcome the limitations caused by most parameters in ZT being closely correlated. Independent 7204

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ACS Nano controls over the ratio between σ and κT (or to be more precise, κE) are not possible because the scattering centers that reduce κT also affect σ. It is also difficult to control S2σ (or more precisely, S2n) because a higher doping concentration (n) will eventually deteriorate S. Through the nanostructuring of spinodally decomposed NPs, we were able to greatly decrease κL and increase μ despite the continual increase in S. We also believe that ensuring the stability of the TE material at higher temperatures (by mixing it with other heterogeneous TE materials and so forth) can result in higher ZT values based on our strategy.

samples was measured with a Netzsch LFA 457 based on the laserflash method. Disks (diameter = 10 mm) of graphite-coated samples were heated from 325 to 623 K. The relative density was calculated from geometric measurements and the mass of the sample, and it was determined to be approximately 96%. A specific heat capacity (Cp) was estimated by using literature values in conjunction with the law of mixtures43 and measured using differential scanning calorimetry (Netzsch DSC 404F3, Supporting Information Figure S7). The total thermal conductivity was calculated from the aforementioned values with the equation κT = DCpd (κT = total thermal conductivity, D = measured thermal diffusivity coefficient, Cp = specific heat capacity, d = density). The carrier densities were estimated with the van der Pauw method in a magnetic field of up to ±2 T.

EXPERIMENTAL SECTION

ASSOCIATED CONTENT

Synthesis. All chemicals were purchased from Sigma-Aldrich (Milwaukee, WI). Phase-separated lead selenide/telluride nanoparticles (PbTe/PbSe NPs) were synthesized according to the following procedures. All experiments were performed in a dry Ar environment. In a typical synthesis, a 3-neck flask containing 7.5 mmol of PbO (99.99%), 2.2 mL of octanoic acid (>99%), 30 mL of diphenyl ether (99%), and 100 mL of trioctylphosphine (TOP, 97%) was heated to 100 °C and degassed for 1−2 h. Then, the temperature was slowly raised to 175 °C while vigorously stirred. Once this temperature was reached, 3 mL of a 0.5 M stock solution of TOP-Se and 4 mL of a 0.5 M stock solution of tributylphosphine (TBP, 97%)-Te,8 which were prepared in a glovebox, were rapidly and sequentially injected into the mixture. Once the TOP-Se solution was injected, the reaction was allowed to proceed for 3 min. The TBP-Te solution was then introduced into the mixture. After 2 min, the reaction was terminated by quenching the mixture to room temperature within 1 min in an ice bath. For the synthesis of pure PbSe NPs, only the TOP-Se solution was injected into the mixture, and the reaction was allowed to proceed for 5 min, with the other procedures the same as those used for the PbSe/PbTe NPs. Both types of NP were dispersed in a mixture of toluene and methanol by performing centrifugation several times. The collected NPs were dried under vacuum conditions at room temperature. Spark Plasma Sintering. To eliminate any remaining surfactants, a heat treatment was performed under low-pressure dry Ar at 300 °C for 30 min. No obvious agglomeration or phase transformation of the NPs occurred, which was confirmed by TEM and XRD analyses. The heat-treated powders were densified with spark plasma sintering (SPS; SPS-20, Welltech) at 350 °C with 40−60 MPa of uniaxial pressure. To minimize grain growth during the sintering process, the samples were cooled without temperature holding. Characterization. X-ray diffraction (XRD; Rigaku D/MAX-2500 V/PC, Japan) and inductively coupled plasma-atomic emission spectrometry (ICP-AES; Varian 710-ES, Australia) were employed to confirm the atomic ratios of Pb, Se, and Te. The overall morphology and crystallographic information were observed with high-resolution field-emission transmission electron microscopy (FETEM; FEI Tecnai G2 F30, 300 kV). To identify the phase transitions, in situ high-temperature XRD (PANalytical X’Pert3 Powder) and in situ TEM (FEI Cryo Technai G2 F20, 200 kV) were used. A focused ion beam system (FIB; JIB-4601F, JEOL) was used to prepare the TEM samples of the sintered nanobulk composites. Thermoelectric Measurements. The sintered pellets were cut and polished for the thermoelectric measurements. The electrical transport and Seebeck coefficient were determined with an ULVAC ZEM-3 system. The prepared samples, which were shaped as rectangular parallelepipeds, were placed between Ni electrodes with two probe thermocouples in contact with one side. To protect the sample surfaces and electrodes, four pieces of graphite foil were inserted between them (we had already confirmed that the graphite foil made no difference to the readings). The sample chamber was heated from 325 to 623 K under approximately 0.1 atm of He. To determine the Seebeck coefficient, the temperature differentials between the hot and cold sides were 10, 20, and 30 °C at each furnace temperature. The thermal diffusivity coefficient (D) of the

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.6b03696. Supporting data and figures (PDF)

AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS The authors thank Dr. H.S.B. of Korea Basic Science Institute (KBSI) for providing the access of their HRTEM. This research was supported by the National Research Foundation (NRF) of Korea grants funded by the Korean government (NRF2013R1A1A2072509). M.-S.K. synthesized the samples, designed and carried out experiments, and wrote the paper. W.J.L. and K.-H.C. helped with sample synthesis and K.-H.C drew some figures. J.-P.A. performed the in situ high-temperature TEM and XRD. Y.-M.S initiated this research by NRF funding and supervised the progress as a principal investigator. M.-S.K., W.-J.L., K.-H.C., J.-P.A. and Y.-M.S conceived the experiments, and analyzed and discussed results. M.-S.K., J.-P.A., and Y.-M.S coedited the manuscript. REFERENCES (1) Baxter, J.; Bian, Z.; Chen, G.; Danielson, D.; Dresselhaus, M. S.; Fedorov, A. G.; Fisher, T. S.; Jones, C. W.; Maginn, E.; Kortshagen, U.; Manthiram, A.; Nozik, A.; Rolison, D. R.; Sands, T.; Shi, L.; Shollh, D.; Wu, Y. Nanoscale Design to Enable the Revolution in Renewable Energy. Energy Environ. Sci. 2009, 2, 559−588. (2) Sales, B. C.; Mandrus, D.; Williams, R. K. Filled Skutterudite Antimonides: A New Class of Thermoelectric Materials. Science 1996, 272, 1325−1328. (3) Hsu, K. F.; Loo, S.; Guo, F.; Chen, W.; Dyck, J. S.; Uher, C.; Hogan, T.; Polychroniadis, E. K.; Kanatzidis, M. G. Cubic AgPbmSbTe2+m: Bulk Thermoelectric Materials with High Figure of Merit. Science 2004, 303, 818−821. (4) Snyder, G. J.; Christensen, M.; Nishibori, E.; Caillat, T.; Iversen, B. B. Disordered Zinc in Zn4Sb3 with Phonon-Glass and ElectronCrystal Thermoelectric Properties. Nat. Mater. 2004, 3, 458−463. (5) Rhyee, J.-S.; Lee, K. H.; Lee, S. M.; Cho, E.; Kim, S. I.; Lee, E.; Kwon, Y. S.; Shim, J. H.; Kotliar, G. Peierls Distortion as a Route to High Thermoelectric Performance in In4Se3‑δ Crystals. Nature 2009, 459, 965−968. (6) Sootsman, J. R.; Chung, D. Y.; Kanatzidis, M. G. New and Old Concepts in Thermoelectric Materials. Angew. Chem., Int. Ed. 2009, 48, 8616−8639. 7205

DOI: 10.1021/acsnano.6b03696 ACS Nano 2016, 10, 7197−7207

Article

ACS Nano (7) Kleinke, H. New Bulk Materials for Thermoelectric Power Generation: Clathrates and Complex Antimonides. Chem. Mater. 2010, 22, 604−611. (8) Wang, X. W.; Lee, H.; Lan, Y. C.; Zhu, G. H.; Joshi, G.; Wang, D. Z.; Yang, J.; Muto, A. J.; Tang, M. Y.; Klatsky, J.; Song, S.; Dresselhaus, M. S.; Chen, G.; Ren, Z. F. Enhanced Thermoelectric Figure of Merit in Nanostructured n-type Silicon Germanium Bulk Alloy. Appl. Phys. Lett. 2008, 93, 193121. (9) Poudel, B.; Hao, Q.; Ma, Y.; Lan, Y.; Minnich, A.; Yu, B.; Yan, X.; Wang, D.; Muto, A.; Vashaee, D.; Chen, X.; Liu, J.; Dresselhaus, M. S.; Chen, G.; Ren, Z. High-Thermoelectric Performance of Nanostructured Bismuth Antimony Telluride Bulk Alloys. Science 2008, 320, 634−638. (10) Ma, Y.; Hao, Q.; Poudel, B.; Lan, Y.; Yu, B.; Wang, D.; Chen, G.; Ren, Z. Enhanced Thermoelectric Figure-of-Merit in p-type Nanostructured Bismuth Antimony Tellurium Alloys Made from Elemental Chunks. Nano Lett. 2008, 8, 2580−2584. (11) Pei, Y.; Shi, X.; LaLonde, A.; Wang, H.; Chen, L.; Snyder, G. J. Convergence of Electronic Bands for High Performance Bulk Thermoelectrics. Nature 2011, 473, 66−69. (12) Pei, Y.; Wang, H.; Snyder, G. J. Band Engineering of Thermoelectric Materials. Adv. Mater. 2012, 24, 6125−6135. (13) Heremans, J. P.; Jovovic, V.; Toberer, E. S.; Saramat, A.; Kurosaki, K.; Charoenphakdee, A.; Yamanaka, S.; Snyder, G. J. Enhancement of Thermoelectric Efficiency in PbTe by Distortion of the Electronic Density of States. Science 2008, 321, 554−557. (14) Heremans, J. P.; Wiendlocha, B.; Chamoire, A. M. Resonant Levels in Bulk Thermoelectric Semiconductors. Energy Environ. Sci. 2012, 5, 5510−5530. (15) Heremans, J. P.; Thrush, C. M.; Morelli, D. T. Thermopower Enhancement in Lead Telluride Nanostructures. Phys. Rev. B: Condens. Matter Mater. Phys. 2004, 70, 115334. (16) Narducci, D.; Selezneva, E.; Cerofolini, G.; Frabboni, S.; Ottaviani, G. Impact of Energy Filtering and Carrier Localization on the Thermoelectric Properties of Granular Semiconductors. J. Solid State Chem. 2012, 193, 19−25. (17) Soni, A.; Yanyuan, Z.; Ligen, Y.; Aik, M. K. K.; Dresselhaus, M. S.; Xiong, Q. Enhanced Thermoelectric Properties of Solution Grown Bi2Te3‑xSex Nanoplatelet Composites. Nano Lett. 2012, 12, 1203− 1209. (18) Zebarjadi, M.; Esfarjani, K.; Dresselhaus, M. S.; Ren, Z. F.; Chen, G. Perspectives on Thermoelectrics: from Fundamentals to Device Applications. Energy Environ. Sci. 2012, 5, 5147−5162. (19) Rowe, D. M. General Principles and Basic Considerations. In Thermoelectric Handbook: Macro to Nano; Rowe, D. M., Ed.; Taylor and Francis: FL, 2006; Ch 1. (20) Biswas, K.; He, J.; Blum, I. D.; Wu, C.-I.; Hogan, T. P.; Seidman, D. N.; Dravid, V. P.; Kanatzidis, M. G. High-Performance Bulk Thermoelectrics with All−Scale Hierarchical Architectures. Nature 2012, 489, 414−418. (21) Zhao, L. D.; Wu, H. J.; Hao, S. Q.; Wu, C. I.; Zhou, X. Y.; Biswas, K.; He, J. Q.; Hogan, T. P.; Uher, C.; Wolverton, C.; Dravid, V. P.; Kanatzidis, M. G. All-Scale Hierarchical Thermoelectrics: MgTe in PbTe Facilitates Valence Band Convergence and Suppresses Bipolar Thermal Transport for High Performance. Energy Environ. Sci. 2013, 6, 3346−3355. (22) Wu, H. J.; Zhao, L.-D.; Zheng, F. S.; Wu, D.; Pei, Y. L.; Tong, X.; Kanatzidis, M. G.; He, J. Q. Broad Temperature Plateau for Thermoelectric Figure of Merit ZT > 2 in Phase-Separated PbTe0.7S0.3. Nat. Commun. 2014, 5, 4515. (23) Wang, H.; Bahk, J.-H.; Kang, C.; Hwang, J.; Kim, K.; Kim, J.; Burke, P.; Bowers, J. E.; Gossard, A. C.; Shakouri, A.; Kim, W. Right Sizes of Nano- and Microstructures for High−Performance and Rigid Bulk Thermoelectrics. Proc. Natl. Acad. Sci. U. S. A. 2014, 111, 10949− 10954. (24) Wang, H.; Hwang, J.; Snedaker, M. L.; Kim, I.-H.; Kang, C.; Kim, J.; Stucky, G. D.; Bowers, J.; Kim, W. High Thermoelectric Performance of a Heterogeneous PbTe Nanocomposite. Chem. Mater. 2015, 27, 944−949.

(25) Zhao, L.-D.; Lo, S.-H.; Zhang, Y.; Sun, H.; Tan, G.; Uher, C.; Wolverton, C.; Dravid, V. P.; Kanatzidis, M. G. Ultralow Thermal Conductivity and High Thermoelectric Figure of Merit in SnSe Crystals. Nature 2014, 508, 373−377. (26) Snyder, G. J.; Toberer, E. S. Complex Thermoelectric Materials. Nat. Mater. 2008, 7, 105−114. (27) Zhao, L.-D.; Dravid, V. P.; Kanatzidis, M. G. The Panoscopic Approach to High Performance Thermoelectrics. Energy Environ. Sci. 2014, 7, 251−268. (28) Kim, M.-S.; Sung, Y.-M. Successive Solution-Liquid-Solid (SLS) Growth of Heterogeneous Nanowires. Chem. Mater. 2013, 25, 4156− 4164. (29) Burch, D.; Bazant, M. Z. Size-Dependent Spinodal and Miscibility Gaps for Intercalation in Nanoparticles. Nano Lett. 2009, 9, 3795−3800. (30) Androulakis, J.; Lin, C.-H.; Kong, H.-J.; Uher, C.; Wu, C.-I.; Hogan, T.; Cook, B. A.; Caillat, T.; Paraskevopoulos, K. M.; Kanatzidis, M. G. Spinodal Decomposition and Nucleation and Growth as a Means to Bulk Nanostructured Thermoelectrics: Enhanced Performance in Pb1‑x SnxTe-PbS. J. Am. Chem. Soc. 2007, 129, 9780−9788. (31) Schwall, M.; Balke, B. Phase Separation as a Key to a Thermoelectric High Efficiency. Phys. Chem. Chem. Phys. 2013, 15, 1868−1872. (32) Poudeu, P. F. P.; D’Angelo, J.; Downey, A. D.; Short, J. L.; Hogan, T. P.; Kanatzidis, M. G. High Thermoelectric Figure of Merit and Nanostructuring in Bulk p-type Na1‑xPbmSbyTem+2. Angew. Chem., Int. Ed. 2006, 45, 1−5. (33) Ibáñez, M.; Zamani, R.; Gorsse, S.; Fan, J.; Ortega, S.; Cadavid, D.; Morante, J. R.; Arbiol, J.; Cabot, A. Core-Shell Nanoparticles as Building Blocks for the Bottom-Up Production of Functional Nanocomposites: PbTe-PbS Thermoelectric Properties. ACS Nano 2013, 7, 2573−2586. (34) Kim, M.-S.; Ryu, J.-J.; Sung, Y.-M. One-Step Approach for Nano-Crystalline Hydroxyapatite Coating on Titanium via Micro-Arc Oxidation. Electrochem. Commun. 2007, 9, 1886−1891. (35) Tan, G.; Shi, F.; Doak, J. W.; Sun, H.; Zhao, L.-D.; Wang, P.; Uher, C.; Wolverton, C.; Dravid, V. P.; Kanatzidis, M. G. Extraordinary Role of Hg in Enhancing the Thermoelectric Performance of p-type SnTe. Energy Environ. Sci. 2015, 8, 267−277. (36) Tian, Z. T.; Garg, J.; Esfarjani, K.; Shiga, T.; Shiomi, J.; Chen, G. Phonon Conduction in PbSe, PbTe, and PbTe1‑xS x from firstprinciples calculations. Phys. Rev. B: Condens. Matter Mater. Phys. 2012, 85, 184303. (37) Pei, Y.; Heinz, N. A.; LaLonde, A.; Snyder, G. J. Combination of Large Nanostructures and Complex Band Structure for High Performance Thermoelectric Lead Telluride. Energy Environ. Sci. 2011, 4, 3640−3645. (38) Jeng, M.-S.; Yang, R.; Song, D.; Chen, G. Modeling the Thermal Conductivity and Phonon Transport in Nanoparticle Composites using Monte Carlo Simulation. J. Heat Transfer 2008, 130, 042410. (39) Sootsman, J. R.; Kong, H.; Uher, C.; D’Angelo, J. J.; Wu, C.-I.; Hogan, T. P.; Caillat, T.; Kanatzidis, M. G. Large Enhancements in the Thermoelectric Power Factor of Bulk PbTe at High Temperature by Synergistic Nanostructuring. Angew. Chem., Int. Ed. 2008, 47, 8618− 8622. (40) Yu, B.; Zebarjadi, M.; Wang, H.; Lukas, K.; Wang, H.; Wang, D.; Opeil, C.; Dresselhaus, M.; Chen, G.; Ren, Z. Enhancement of Thermoelectric Properties by Modulation-Doping in Silicon Germanium Alloy Nanocomposites. Nano Lett. 2012, 12, 2077−2082. (41) Pei, Y.-L.; Wu, H.; Wu, D.; Zheng, F.; He, J. High Thermoelectric Performance Realized in a BiCuSeO System by Improving Carrier Mobility through 3D Modulation Doping. J. Am. Chem. Soc. 2014, 136, 13902−13908. (42) Non-Tetrahedrally Bonded Elements and Binary Compounds I, Landolt-BörnsteinGroup III Condensed Matter; Madelung, O., Rössler, U., Schulz, M., Eds.; Springer: New York, 1998. 7206

DOI: 10.1021/acsnano.6b03696 ACS Nano 2016, 10, 7197−7207

Article

ACS Nano (43) Blachnik, R.; Igel, R. Thermodynamic Properties of IV-VI Compounds Lead Chalcogenides. Z. Naturforsch., B: J. Chem. Sci. 1974, 29, 625−629. (44) Heremans, J. P.; Thrush, C. M.; Morelli, D. T. Thermopower Enhancement in PbTe with Pb Precipitates. J. Appl. Phys. 2005, 3, 063703. (45) Wang, H.; Schechtel, E.; Pei, Y.; Snyder, G. J. High Thermoelectric Efficiency of n-type PbS. Adv. Energy Mater. 2013, 3, 488−495. (46) Wu, D.; Zhao, L.-D.; Tong, X.; Li, W.; Wu, L. J.; Tan, Q.; Pei, Y. L.; Huang, L.; Li, J.-F.; Zhu, Y.; Kanatzidis, M. G.; He, J. Superior Thermoelectric Performance in PbTe-PbS Pseudo-Binary: Extremely Low Thermal Conductivity and Modulated Carrier Concentration. Energy Environ. Sci. 2015, 8, 2056−2068. (47) Gelbstein, Y.; Davidow, J.; Girard, S. N.; Chung, D. Y.; Kanatzidis, M. Controlling Metallurgical Phase Separation Reactions of the Ge0.87Pb0.13Te Alloy for High Thermoelectric Performance. Adv. Energy Mater. 2013, 3, 815−820. (48) Doak, J. W.; Wolverton, C. Coherent and Incoherent Phase Stabilities of Thermoelectric Rocksalt IV-VI Semiconductor Alloys. Phys. Rev. B: Condens. Matter Mater. Phys. 2012, 86, 144202. (49) Day, T. W.; Zeier, W. G.; Brown, D. R.; Melot, B. C.; Snyder, G. J. Determining Conductivity and Mobility Values of Individual Components in Multiphase Composite Cu1.97Ag 0.03Se. Appl. Phys. Lett. 2014, 105, 172103. (50) Gelbstein, Y. Morphological Effects on the Electronic Transport Properties of Three-Phase Thermoelectric Materials. J. Appl. Phys. 2012, 112, 113721. (51) Gelbstein, Y. Phase Morphology Effects on the Thermoelectric Properties of Pb0.25Sn0.25Ge0.5Te. Acta Mater. 2013, 61, 1499−1507. (52) Wu, D.; Zhao, L.-D.; Hao, S.; Jiang, Q.; Zheng, F.; Doak, J. W.; Wu, H.; Chi, H.; Gelbstein, Y.; Uher, C.; Wolverton, C.; Kanatzidis, M.; He, J. Origin of the High Performance in GeTe-Based Thermoelectric Materials upon Bi2Te3 Doping. J. Am. Chem. Soc. 2014, 136, 11412−11419. (53) Hazan, E.; Ben-Yehuda, O.; Madar, N.; Gelbstein, Y. Functional Graded Germanium-Lead Chalcogenide-Based Thermoelectric Module for Renewable Energy Applications. Adv. Energy Mater. 2015, 5, 1500272. (54) Chen, N.; Gascoin, F.; Snyder, G. J.; Müller, E.; Karpinski, G.; Stiewe, C. Macroscopic Thermoelectric Inhomogeneities in (AgSbTe2)x(PbTe)1‑x. Appl. Phys. Lett. 2005, 87, 171903. (55) Li, J.-F.; Liu, W.-S.; Zhao, L.-D.; Zhou, M. High-Performance Nanostructured Thermoelectric Materials. NPG Asia Mater. 2010, 2, 152−158.

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