Significant Enhancement of the Thermoelectric Performance of Higher

Article pubs.acs.org/cm

Significant Enhancement of the Thermoelectric Performance of Higher Manganese Silicide by Incorporating MnTe Nanophase Derived from Te Nanowire Zhiliang Li,† Jin-Feng Dong,† Fu-Hua Sun,† Shinsuke Hirono,‡ and Jing-Feng Li*,† †

State Key Laboratory of New Ceramics and Fine Processing, School of Materials Science and Engineering, Tsinghua University, Beijing 100084, P. R. China ‡ Department of Advanced Materials, Toyota Automobile Technology Co., Ltd., Susono, Shizuoka 410-1193, Japan S Supporting Information *

ABSTRACT: Higher manganese silicide (HMS) is a naturally abundant, eco-friendly, and low-cost p-type thermoelectric semiconductor with high power factor (PF); however, its figure of merit (ZT) is very low due to an intrinsically high thermal conductivity (κ). It is also difficult to synthesize a HMS single phase without MnSi or Si second phase, which generally leads to a significant decrease of the PF and/or increased κ. In this study, pure HMS was obtained via wet ball milling in combination with spark plasma sintering. A further p-type MnTe, with high Seebeck coefficient (S) and low κ, was incorporated into the HMS matrix to form MnTe/HMS composites. To effectively decrease κ, while simultaneously enhancing the electrical conductivity, Te nanowires of ∼35 nm in diameter were synthesized with a solution phase method and subsequently embedded to HMS to form MnTe/ HMS nano/bulk structures. The incorporation of Te nanowires led to a 38% reduction of κ in addition to a slight increase of S. Finally, the ZTmax of MnTe/HMS nano/bulk composites increased by ∼71% from 0.41 to 0.70. This study demonstrates a unique and facile method to boost the ZT of HMS to a high level, which is also applicable to other thermoelectric materials.

■

INTRODUCTION Thermoelectric (TE) materials have attracted increasing attention over the past few decades not only due to their excellent abilities for directly converting electricity from temperature gradients, but also due to their broad application prospects in solid-state electronic refrigeration.1−4 Generally, the dimensionless figure of merit (ZT) governs the TE energy conversion efficiency. ZT is defined as ZT = S2σT/κtot, where S, σ, T, and κtot are the Seebeck coefficient, electrical conductivity, absolute temperature, and total thermal conductivity, respectively.2,5−8 Currently, the best commercially available and most efficient TE materials at medium temperatures are PbTe based compounds9,10 and SnSe single crystals11,12 with maximum ZT values of about 2.2 and 2.6, respectively. Nevertheless, numerous difficult tasks must be overcome before their broad commercial application can be realized. These tasks are related to the current complicated preparation processes, the utilization of the toxic Pb element, as well as thermal and chemical instabilities. In contrast, among midtemperature TE materials, higher manganese silicide (HMS) is a naturally abundant, ecofriendly, and low-cost TE semiconductor with satisfactory thermal stability due to high melting point and good mechanical strength. As a p-type semiconductor, HMS has a band gap of about 0.77 eV, and the electronic density of the © 2017 American Chemical Society

state near the Femi energy changes sharply, which usually causes a high Seebeck coefficient (S = 130−220 μVK−1) and a good electrical transport property (σ = 70−33 kSm−1).13,14 Thus, the power factor (PF = S2σ) of pure HMS ranged from 0.7 × 10−3 to 1.5 × 10−3 Wm−1 K−2, which is even comparable to the PF of unmodified PbTe0.7S0.315 or Ge0.87Pb0.13Te16 alloys, whose ZT values are higher than 2.0 after some modifications. This suggests HMS to be a particularly promising TE material. However, state-of-the-art research has revealed the ZT values of pure HMS to be only about 0.2−0.4 at 773 K. The intrinsically high total thermal conductivity (κtot) of 3.0 ± 0.2 Wm−1 K−1 at 773 K is about eight-times higher than that of monocrystal SnSe (0.25−0.35 Wm−1 K−1, from 673−973 K),11 which has become a key limiting issue to the TE performance of HMS. Currently, significant efforts have been made to reduce the κtot of HMS. Truong et al.17 reported that when the average grain size of HMS was reduced from 100 to 10 μm, the κtot decreased by 27% (from 3.3 to 2.4 Wm−1 K−1 at 773 K), and the corresponding ZT increased from 0.28 to 0.54; Shi et al.18,19 demonstrated that the κtot can be decreased to 2.2 Wm−1 Received: June 3, 2017 Revised: July 29, 2017 Published: August 7, 2017 7378

DOI: 10.1021/acs.chemmater.7b02270 Chem. Mater. 2017, 29, 7378−7389

Article

Chemistry of Materials

Figure 1. (A) Unit cell information on HMS phases; (B) XRD patterns with different raw ratios of Mn/Si, ranging from 1:1.70 to 1:1.81; (C) lowmagnification, (D) magnified, and (E) HRTEM images of HMS; (F) simulative atomic arrays of HMS.

K−1 at 773 K (a decrease of about 16%) when ReSi1.75 precipitates of 50−200 nm in diameter formed in the Re substituted HMS samples, and consequently, the ZT was enhanced from 0.45 to 0.57 (increased by 27%); Takeuchi et al.20 proved that the κtot can be further decreased to 1.8 Wm−1 K−1 at 773 K (decreased by 33%) when 6 at% Re was used to form a supersaturated Re solid solution HMS, and the corresponding ZT value was increased to 0.9 at 773 K (an increase of about 87%, the highest ZT in this report is 1.04 at 923 K). These phenomena suggest that most of the contribution of the enhanced ZT resulted from the thermal conductivity and that κtot plays a decisive role in improving the TE properties of HMS. In fact, κtot principally consists of electrical thermal conductivity (κe) and lattice thermal conductivity (κL); in many semiconductors, κL is typically much higher than κe, which is normally restricted by poor

electrical conductivity (κe = LσT; L and σ are Lorentz number and electrical conductivity).5 Most efforts to reduce the κL of HMS focused on purification to remove the impurity monatomic Si or semimetallic MnSi phase,21−23 processing element doping to increase lattice defects,14,24−28 intensifying grain refining to increase the amount of grain boundaries,17,29−31 and precipitating nanoinclusions to intensify the phonon scattering.18,32−38 Consequently, the ZT values could be enhanced to 0.43−0.55, 0.57−0.65, 0.52−0.62, and 0.43− 1.04, respectively. Nevertheless, many limitations still remain that should be resolved to further improve the TE properties. First, some facile, green, rapid, and energy-saving preparation technologies such as mechanical alloying, spark plasma sintering (SPS) rather than arc melting, melt spinning, or chemical vapor transport methods are required to be applied or need to be improved to obtain pure HMS with refined grains. 7379


Article

Chemistry of Materials Second, while reducing κL, sometimes the electrical transport properties are adversely affected, and this should be avoided as much as possible. Third, more studies on how to precipitate desired inclusions as designed are urgently necessary because size, shape, and composition of the inclusions as well as their dispersion are generally uncontrollable in this method. Moreover, more investigations are needed to study whether the predicted size effects derived from the inclusions can be achieved in experiments. Thus, in this study, we prepared pure HMS by adjusting the ratio of Mn/Si via a simple, rapid, and low-energy consumption avenue: wet ball milling combined with spark plasma sintering. To avoid a potentially negative effect on the electrical transport properties caused by the doping effect, a p-type TE semiconductor (MnTe) with good thermal and chemical stabilities was selected to form MnTe/HMS composite structures. To further decrease the κL of MnTe/HMS composite, Te nanowires with a diameter of about 35 nm and a length over 8 μm were synthesized via chemical solution phase method and subsequently embedded into the HMS to form MnTe/HMS nano/bulk structures. A surface modification was also performed on the Te nanostructures using a mixture of anhydrous ethanol and normal hexane to obtain well-dispersed nano/bulk composites. The final results show that κ tot decreased by about 38% in the 3.0% MnTe/HMS nano/bulk specimen, and the corresponding ZT increased by ∼71% from 0.41 to 0.70 at 823 K, which resulted in one of the highest values among the state-of-the-art findings on HMS.

obtained. The corresponding refined calculation reveals lattice constants a and c of 5.525 and 65.550 Å, respectively, which correspond well to the JCPDS card No. 72−32 (a = b = 5.531 Å, c = 65.311 Å). Although the element ratio of Mn/Si was 1:1.80, the final sample was dominantly composed of Mn15Si26 (MnSi1.733), which means that ∼0.37% of Si element was lost in wet ball milling due to the relatively strong stickiness between the Si element and balls as revealed by Sadia et al.43 However, when the element ratio of Mn/Si was further reduced to 1:1.81, the Bragg peaks of Si crystal (JCPDS card No. 89−5012), (111) plane (2θ = 28.5°) and (311) plane (2θ = 56.1°) arose. Thus, the relatively narrow homogeneity range of HMS (Mn/Si from 1:1.79 to 1:1.81) and the emulative and alternative features47 between MnSi phase and Si crystal have been proven with this method. Low-magnification, high-resolution transmission electron microscopy (HRTEM), selected area electron diffraction (SAED) pattern, and simulative lattice orientation arrays are shown in Figure 1C−F. The final pure HMS bulk consisted of single crystals with diameter variations from 400−900 nm, and numerous nanopores or defects could also be observed in the internal of crystals as well as on the grain boundaries, which probably stem from the intrinsic characteristic of wet ball milling. Some parallel lattice fringes are shown in Figure 1D, and the corresponding SAED pattern along the [1−10] zone axis is displayed in the inset of Figure 1D. Interestingly, because of the strong superlattice orientation of Mn and Si subsystem lattices among every periodic lattice fringe with dense atom arrays, a series of subordinate spots appeared in both diffraction spot and center transmission spot, which was also reported by Shi et al.19 The shortest interplanar spacings between center diffraction spots and center transmission spots were approximately 0.2363, 0.2469, and 0.2427 nm, which corresponds to the Mn15Si26 (004), (121), and (125) planes, respectively. All these planes belong to the (hkl, l ≠ 0) plane family, indicating that these lattice fringes are perpendicular to the [001] direction of HMS. The detailed atomic arrangement information is plotted in Figure 1E, and similar interplanar spacings for the 0.2325 nm (004) plane, 0.2452 nm (121) plane, and 0.2418 nm (125) plane were also observed. The corresponding simulative atomic arrays are depicted in Figure 1F, which reveals that the closely pitched atomic arrangements consist of the relatively dense atomic configuration of the HMS unit cell. This phenomenon again proved that the lattice fringes are perpendicular to the [001] direction of the HMS. Generally, trace amounts of monatomic Si or semimetallic MnSi easily separate out in the HMS matrix due to the relatively narrow homogeneity range; thus, it is necessary to evaluate how much influence different types of inclusions will exert on the TE properties. Detailed TE coefficients are plotted in Figure S1. For the pure HMS, the PFmax is about 1.6 × 10−3 Wm−1 K−2 at 623 K, which is even higher than reported in the most current reports. Nevertheless, when Si or MnSi has been precipitated in the HMS, the PFmax respectively decreased to 1.38 × 10−3 Wm−1 K−2 and 1.30 × 10−3 Wm−1 K−2 due to their dramatic decrease on σ or S. The κtot of the pure HMS at 773 K, where generally has the optimal ZT value is about 3.0 Wm−1 K−1, which is similar to the results of numerous reports.19,28,48,49 By taking the high PF into account, the maximum ZT of the pure HMS was about 0.41 at 773 K (red line), which is comparable or marginally higher than reported in most other current reports.18,45,50−53 However, κtot of HMS with MnSi inclusions slightly decreased due to

■

RESULTS AND DISCUSSION HMSs normally crystallize to a Nowotny chimney-ladder structure,39,40 which can be described as an inlay of Mn and Si sublattices, where the Mn atoms form the chimney frame wall, and in which the Si atoms are arranged as spiral ladders. Therefore, in the tetragonal HMS, the lattice parameter a is predominantly restricted by the Mn unit cell, while the c parameter can vary from 17 to 118 Å, which strongly depends on the accumulation situation of Si atoms. In fact, HMS is a general term that includes a family of Nowotny chimney-ladder phases.41,42 A number of different structural formulas of such commensurate phases include Mn4Si7, Mn7Si12, Mn11Si19, Mn15Si26, Mn19Si33, Mn26Si45, Mn27Si47, and Mn39Si68, which have been reported with determined crystal structures (Figure 1A). Fortunately, all of these variations show little variation in their TE properties and remain stable up to above 800 °C. In contrast, because of the ultranarrow homogeneity range of HMS, two side phases, semimetallic MnSi and monatomic Si easily appeared in the HMS matrix. Specifically, the MnSi impurity phase is very difficult to be removed due to the relatively low formation enthalpy and the excellent chemical stability when HMS is fabricated via mechanical alloying,22,43,44 arc melting,32,36,45 solid-state reaction,46 melt spinning,29 and chemical vapor deposition.30 Thus, to avoid the formation of second phases, the facile, rapid, and energy-saved avenue of wet ball milling, combined with SPS, was used to prepare pure HMS. As the XRD pattern (Figure 1B) shows, the Bragg diffraction peaks of the MnSi phase (JCPDS card No. 81−483), (110) plane (2θ = 27.7°), (111) plane (2θ = 34.1°), (210) plane (2θ = 44.5°), and (211) plane (2θ = 49.0°) gradually reduced as the ratio of Mn/Si decreased from 1:1.70 to 1:1.79. The MnSi phase was not detected when the ratio of Mn/Si decreased to 1:1.80 (see the red line in Figure 1B), where virtually single HMS phase was 7380


Article


Figure 2. (A) XRD patterns of MnTe/HMS composites with different nominal ratios of (MnSi1.80)1−x(MnTe)x, x ranging from 0−5.0 mol %; (B) SEM and (C) TEM images of the (MnSi1.80)0.97(MnTe)0.03 sample; (D, E) likely formation mechanisms of the porous MnTe precipitate.

these XRD patterns, no other peak of impurity phase like Si or MnSi phase was detected, which means that the final bulks only consisted of the MnTe and Mn15Si26, as we expected. The corresponding SEM images of the MnSi1.746Te0.03 sample are shown in Figure 2B and inset, some white crystals with diameters in the micrometer range were precipitated along the gray big crystals, and similar phenomena were also found in other specimens with different Te powder ratios. To accurately confirm which ingredients the white precipitates contain, an electronic probe microanalyzer (EPMA) was performed on the MnSi1.71Te0.05 sample. As the EPMA images (Figure S2) indicate, the Mn elements were uniformly distributed, while some small Te enriched regions and numerous big Si enriched areas were observed in Figure S2C,D. Correspondingly, this proved that the white precipitates were MnTe crystals, and the big matrixes were HMS. Their detailed structure information is also provided in the magnified TEM (Figure 2C), which reveals that the micrometer MnTe precipitates consists of nanoparticles with nearly coherent lattices. Furthermore, numerous nanopores are also observed among the MnTe precipitates. The likely formation mechanisms are described in Figure 2D and E, and the corresponding chemical reaction equations are given as follows:

intensified phonon scattering, while the monatomic Si has some negative effect on reducing the κtot due to its intrinsically high κtot (≥60 Wm−1 K−1). As a result, the ZT values of the HMS samples with Si or MnSi inclusions were only 0.29 and 0.38, respectively. All these phenomena elucidate neither Si nor MnSi to be a suitable precipitate to enhance the TE properties of HMS; therefore, more ideal materials with intrinsically low κtot and high Seebeck coefficient are required to reduce the κtot of HMS as well as to maintain the relatively high PF. MnTe is also a p-type of TE material with a relatively broad band gap (∼1.3 eV), and it has the same responsive temperature range (from 300−800 K) to HMS. However, the S of MnTe (from 650−300 μVK−1) is far higher than that of the HMS (from 100−210 μVK−1). Interestingly, in reversal of HMS, MnTe transforms to be metallic when the temperature increased to 673 K and then displayed a drastic σ increase with increasing temperature, which could make a positive contribution for the MnTe/HMS composite at high temperatures. Moreover, the κtot of MnTe is only about 0.55 Wm−1 K−1 at 773 K, which is about five-times lower than that of HMS.54−59 Thus, it is possible to get a relatively low κtot as well as to retain the high electrical transport properties when MnTe is incorporated into HMS. Finally, MnTe has good thermal stability, whose melting point (1150 °C) is close to that of HMS (1140 °C), making it possible to coexist with HMS at 1050 °C (SPS temperature to synthesize HMS). Different Te powder amounts (from 0.5−5.0 mol %) were ad d ed t o t h e m ix t u r e o f M n an d S i t o fo r m (MnSi1.80)1−x(MnTe)x composites. Their XRD patterns are shown in Figure 2A, and lines a and i show the MnTe and Mn15Si26 JCPDS card No. 65−5047 and No. 89−2413, respectively. Lines b to h are MnTe/HMS composites with different amounts of Te, and these are also shown in the inset of Figure 2A, the diffraction peak of MnTe, like (101), (102), and (110) planes, increased with increasing Te amount. In all of

2Te + Mn → MnTe2

(1)

MnTe2 → Te + MnTe

(2)

Te + Mn → MnTe

(3)

when the temperature below 430 °C, the intermediate MnTe2 compound first formed due to its relatively high formation energy (eq 1, Figure S3). Then the Te atoms gradually diffused out with increasing temperature due to the poor thermal and chemical stabilities of MnTe2.60−62 The dissociative Te atoms can react with the near Mn atoms, thus forming new MnTe crystals (eqs 2 and 3 and Figure 2D). Furthermore, many pores 7381


Article


Figure 3. (A) Seebeck coefficients, (B) electrical conductivities, (C) thermal conductivities, and (D) ZT values of (MnSi1.80)1−x(MnTe)x composites as a function of temperature; (E) simple simulation mechanism of the intensified phonon scattering effect.

once filled up with Te atoms were left (Figure 2E), which probably also intensified the phonon scattering and depressed κL. Only marginal variance occurred on the carrier concentrations when 3.0 mol % Te powder was added to the HMS (Figure S4A), which means that the doping effect did not happen in the final sample. Similar to the pure HMS, the carrier mobility of the MnSi1.746Te0.03 specimen using Te powder as a function of 1000T−1 still fit to the T−3/2 curve well instead of the T−1/2 curve (Figure S4B), indicating that the dominant carrier scattering mechanism belonged to acoustic scattering rather than to alloy scattering.63 In addition, as shown in Figure S4C, the lattice parameters a and c had little change with Te element increasing. All these phenomena proved that the final specimens were MnTe/HMS composites instead of solid solutions. Because of the superimposed effect derived from the addition of MnTe, the S at 773 K gradually increased from 212 to 225 μV K−1 with the Te element increasing from 0 to 5.0 mol % (Figure 3A). As a reference, the pure MnTe bulk was also prepared with the same method, and the corresponding S is given in Figure S5 as ∼337 μV K−1 at 773 K, which is far higher than that of HMS. Therefore, the more MnTe particles

precipitated, the higher was the S value. However, depending on the linear effect of 3% MnTe and 97% HMS, the theoretical S was still a little lower than the experimental values. Probably, the enhanced energy filtering effect may be the simplest cause for the S difference.64−69 As theoretical studies demonstrated, the relationship between the relaxation time (τ) and the energy (E) of the carrier can be defined by τ = τ0Eγ−1/2, where τ0 is a constant irrelevant to E, and γ is the scattering parameter. Normally, the Eg of MnTe (1.3 eV) is higher than that of HMS (0.77 eV). Thus, when some MnTe particles were precipitated in the HMS matrix, a complex band gap with many energy barriers and potential wells was formed in the final composites. The holes with low energy usually have shorter relaxation times and can be filtered or trapped more easily than the high-energy carriers by the energy barrier or by potential wells when transmitting. Therefore, only high-energy carriers can be gathered at the ends of the dense samples, which ultimately results in a relatively high S. However, as the Te element increased from 0 to 5.0 mol %, the σ decreased by 17%, from 35.4 (black line) to 29.3 kSm−1 (pink line) at 773 K (Figure 3B) due to the addition of MnTe, with σ of only 14% of HMS. At the same time, the increased porous structures and grain 7382


Article


Figure 4. (A) XRD patterns of Te nanowires; (B) SEM image of Te nanowires; (C) Te nanowires distributed throughout different solvents; (D) SEM images of the mixture of Te nanowires, Si, and Mn powders; (E) dissolution mechanism of Te nanowire in the mixture of alcohol and nhexane.

boundaries also affected the final σ negatively. Ultimately, as a consequence of the much-decreased electrical conductivities, the PF slowly decreased from 1.60 × 10−3 (pure HMS) to 1.49 × 10−3 Wm−1 K−2 with the Te element increasing to 5.0 mol %. Interestingly, some distinct decrease of κ occurred when MnTe was incorporated into HMS. κtot first rapidly reduced from 2.95 to 2.20 Wm−1 K−1 as the Te element increased to 3.0 mol % and then slowly decreased to 1.97 Wm−1 K−1 at 773 K as the Te element ratio continually increased to 5.0 mol % (Figure 3C). The ∼35% decrease from pure HMS to the (MnSi1.80)0.95(MnTe)0.05 sample was principally caused by the κL, whose value was reduced from 2.53 to 1.62 Wm−1 K−1, decreasing by ∼36% as shown in the inset of Figure 3C. Thus, the intensified phonon scattering is the dominant mechanism for reducing κtot of MnTe/HMS composites. Probably, the abundant pores70,71 as well as the increased grain boundaries and lattice defects offered a great contribution to scattering the low-frequency, mid-frequency, and high-frequency phonons, respectively, which enabled few phonons to transmit from the

hot side to the cold side (Figure 3E). Furthermore, because of the relatively high Eg of MnTe, which generally led to a weak bipolar diffusion effect, the κtot value at 823 K, which was once much higher than that of 773 K, decreased more and more, while the Te element ratio gradually increased to 5.0 mol %. Ultimately, as a consequence of the relatively stable PF and continuously decreased κtot, the ZT increased by about 46%, from 0.41 (Pure HMS) to 0.60 ((MnSi1.80)0.97(MnTe)0.03) shown in Figure 3D. MnTe is the right option to enhance the TE performance of HMS. In the preparation of the (MnSi1.80)1−x(MnTe)x composites, Te powder with a diameter above 38 μm was used, and nanometer MnTe precipitates with poor dispersion (the overall diameter of precipitates were in micrometer, as shown in Figure 2B and inset) were formed along the grain boundaries of HMS crystals. In fact, numerous theoretical studies have proved that if the size of the scattering center decreases to nanoscale and is uniformly dispersed, κtot will sharply decrease and little impacts will happen on the σ due to the relatively short mean free path 7383


Article


Figure 5. (A) SEM and (B) TEM images of MnTe/HMS nano/bulk compounds; corresponding EDS on (C) point I and (D) point II in panel B; (E) detailed HRTEM image on the grain boundary of the nano/bulk structure; (F) possible formation mechanism of MnTe inclusions with intermediary grain boundary.

morphologies of the final bulks are shown in Figures 5A and S6. In contrast to (MnSi1.80)1−x(MnTe)x composites whose Te precursor and MnTe inclusions range in the micrometer scale, many MnTe precipitates with diameters below 100 nm were observed in this case. Furthermore, the MnTe nanoinclusions were uniformly dispersed in the HMS bulk (we named it MnTe/HMS nano/bulk structure). Moreover, many nanosize pores, which probably stemmed from the volatilization of residual solvents, were also found in the final bulks. Two particular morphologies of nanoinclusions are presented in Figures 5B and S5C. The diameter of these nanoparticles was 85 ± 10 nm. To further confirm the ingredients of these particles, a contrastive EDS between point I and point II was obtained, which is shown in Figure 5C and D, respectively. At point I, the EDS peak of the Te element was too weak to be observed due to its low atomic ratio (0.02%). Nevertheless, an inverse phenomenon happened at point II, the EDS peaks of Te element remarkably increased with decreasing Si element, and finally, the atomic ratio of Te/Mn increased to 49.12:49.42, which corresponded well to the proportion of MnTe. More interestingly, as indicated by the HRTEM image (Figure 5E), an intermediary grain boundary with a diameter of about 20 nm was formed between MnTe inclusion and HMS matrix. The corresponding possible formation mechanism is also shown in Figure 5F. In the first step, a compact bulk with Te NWs fully surrounded by Mn atoms was obtained under high axial pressure (60 MPa) during SPS. Then a combination reaction between the Te NWs and Mn atoms happened when the SPS temperature stopped at 430 °C for 30 min, whereby the Te NWs transformed into MnTe2 NWs. Furthermore, the diameter of the nanowire increased due to the lattice parameter difference between Te (a = b = 0.445 nm) and MnTe2 (a = b = 0.695 nm). However, MnTe2 NWs were not the final inclusions; as the temperature increased, the excess Te atoms gradually diffused out from the MnTe2 NWs and reacted with nearby Mn atoms to form intermediary grain boundaries (IGBs, right inset of Figure 5F) and correspondingly led to a doubling of the diameter of the NWs. To ensure the influence the Te NWs will exert, TE performance testing was performed in the (MnSi1.80)0.97(MnTe)0.03 nano/bulk composite with Te NWs

of carriers.10,72 Therefore, in the following experiments, Te nanowires (NWs) were synthesized and mixed with a mixture of Si and Mn powder to prepare the MnTe/HMS nano/bulk structures to further reduce the κtot and enhance electrical transport properties. To obtain the uniform Te NWs, a facile and efficient solution phase method was used, and preliminary steps have been reported in our previous publications.73,74 As shown by the XRD patterns (Figure 4A), the diffraction peaks of the final sample corresponded well with the JPCDS card No. 86−2269, and the (100), (101), (012), (110), and the (021) planes of the hexagonal Te were easily discerned. After adjusting the injection rate of the reducing agent (hydrazine hydrate), uniform Te NWs with average diameters of about 35 nm and lengths above 8.0 μm (length/diameter ratio larger than 220:1) were obtained. Furthermore, with the help of a polyvinylpyrrolidone (PVP) surfactant, the final Te NWs were evenly distributed in alcohol as shown in Figure 4B. However, some Te nanowire agglomerations had formed when the Te NWs were directly dispersed in n-hexane (left inset of Figure 4C). Thus, a surface modification of Te NWs was required to obtain well-dispersed Te NWs. To the best of our knowledge, even though the Te NWs were carefully washed several times with deionized water and alcohol, an ultrathin PVP molecular layer with hydrophilic group outside was still capped on the surface of Te NWs (Figure 4E). Thus, the final Te NWs were finely distributed in an aqueous solvent such as alcohol, water, or ethylene glycol rather than an oily solvent such as n-hexane or methylbenzene. On the basis of the collection of test results, we finally used alcohol to bridge PVP and n-hexane. As the molecular dissolution mechanism (Figure 4E) revealed, two mutual dissolutions at both sides of alcohol happened due to the principle of the dissolution in the similar structure. As a result, the Te NWs first evenly dispersed in the alcohol and then dispersed well in n-hexane (right inset of Figure 4C). Ultimately, a uniform mixture of Te NWs, Si, and Mn powders was obtained (Figure 4D and inset). The final mixtures of Te NWs, Si, and Mn powders were sintered to bulks via SPS. To avoid Te NWs degeneration to micro scale, presintering was performed at 430 °C for 30 min to first transform the Te NWs into MnTe2 NWs. The detailed 7384


Article


Figure 6. (A) Seebeck coefficients, (B) electrical conductivities, (D) power factors, (E) total thermal conductivities, (F) lattice thermal conductivities, and (H) ZT values of MnTe/HMS compounds; (C) carrier concentration and mobility of different samples; (G) 3D simulated mechanism of phonon scattering on the MnTe/HMS nano/bulk structure.

723 K (Figure 6D). These values are even higher than those of pure HMS (1.60 × 10−3 Wm−1 K−2) as well as those reported in state-of-the-art publications.18,22,29,32 Moreover, the relatively uniform nano/bulk structure may also affect κtot. Although Te NWs were used in the sample with n-hexane as solvent, κtot seems a little higher than that MnSi1.746Te0.03 specimen using Te powder, and the Te nanowire agglomerations may contribute a relatively weak nanometer size effect to the phonon scattering. When a mixture of alcohol and n-hexane was used to evenly disperse the Te NWs, κtot further decreased to 1.84 Wm−1K−1 at 773 K (Figure 6E), which was 62.3% for pure HMS or 91.1% for the MnSi1.746Te0.03 sample using Te powder. The κL of these samples can be roughly estimated via κL = κtot − κe = κtot − LσT, in which L is ∼1.54 × 10−8 V2 K−2 for pure HMS.18 The corresponding κL values are described in Figure 6F, where a similar variation trend was observed for κtot. To reveal how low κL is, a theoretical κL min of pure HMS was provided by the follow equations, as proposed by Cahill et al.:75

as precursor. The influence that stemmed from different dispersants was also studied. When Te NWs were used, regardless of the matter distribution in the n-hexane or the mixture of alcohol and n-hexane (Valcohol/Vn‑hexane = 1:5), the S at 773 K (blue and green lines in Figure 6A) was marginally higher than that of the sample with Te powder as raw material. We conjecture that the energy filtering effect64−69 was further intensified by the increased amount of scattering centers when the Te micrometer powder was replaced by Te NWs. This assumption was proven by the continually reduced carrier concentrations shown in Figure 6C (red bars). Generally, when the energy filtering effect happened, effective carrier concentrations decreased due to the enhanced carrier scattering derived from the inclusions. Furthermore, as revealed by Figure 6B, in the whole temperature range, the electrical conductivities seemed to increase when Te NWs were used. Particularly, when the Te NWs were evenly dispersed in the mixture of alcohol and n-hexane, uniform nano/bulk structures were obtained, which ultimately offered a relatively unblocked and compact passageway for the carriers. Therefore, although the carrier concentrations decreased slightly, the acutely increased mobility still led to a higher σ in the MnSi1.746Te0.03 specimen using Te NWs than that using Te powder. Ultimately, considering the relatively high S and σ, the maximum PF of the sample using Te NWs further increased from 1.56 × 10−3 (MnSi1.746Te0.03 specimen using Te powder) to 1.73 × 10−3 Wm−1 K−2 at

κLmin =

⎛ T ⎞2 ⎛ π ⎞1/3 2/3 ⎜ ⎟ κBn ∑ νi⎜ ⎟ ⎝6⎠ ⎝ θi ⎠ i

∫0

θi / T

x 3e x dx (e − 1)2 x

(4)

where κB, n, vi, and θi are Boltzmann constant, the number density of atoms, speeds of sound,18 and the cutoff frequency for each polarization, respectively. As Figure 6F indicates, when 7385


Article


60 °C for 4 h. Then a two-step sintering process (430 °C for 30 min, and 1050 °C for 6 min) was performed in spark plasma sintering (SPS211Lx, Fuji Electronic Industrial Co., Ltd., Japan) under an axial pressure of 60 MPa and a vacuum of 3.0 Pa. Finally, the MnTe/HMS bulk specimen with diameter ∼15 mm was obtained, and a disk with thickness about 1.5 mm and diameter about 10 mm was cut from the direction perpendicular to pressing force to evaluate the thermal conductivity. A rectangular column with a size of 7 × 2.5 × 2.5 mm3 was sliced out along the direction parallel to axial pressure to measure the electrical transport properties. The MnTe/HMS composites with nano/bulk structures were obtained by replacing the Te powder with Te nanowires in similar reaction conditions. Controlled Synthesis of Te NWs. Te NWs were synthesized via facile solution phase method, which has been reported in our previous studies.73,74 PVP (3.6 g) and 18 mmol TeO2 were added to 300 mL of ethylene glycol (EG) in a 500 mL three-necked flask with vigorous and continuous magnetic stirring. The mixture was heated to 170 °C and then turned to gray when 1.0 mL of hydrazine hydrate (N2H4·H2O) was evenly injected into the solution. After continual stirring for 45 min, immaculate Te NWs were obtained. The final colloids were left to naturally cool to room temperature. The Te NWs were collected via centrifugation at 4000 rpm, washed several times with absolute ethanol and distilled water, and adequately dried to powder. The final Te NWs were distributed into the intermixture of n-hexane and anhydrous ethanol and ball milled with the Si and Mn powders to prepare a homogeneous elementary substance mixture. Characterization Methods. The crystal phases of the specimens were studied using a Bruker AXS XRD-D8 Focus X-ray diffractometer (XRD, Germany), equipped with graphite monochromatized CuKα radiation (λ = 1.5406 Å). A field emission scanning electron microscope (FESEM, Zeiss Merlin, Germany) was used to characterize the morphologies of the as-prepared products. High-resolution transmission electron microscopy (HRTEM) images and selected area electron diffraction (SAED) patterns were recorded via JEOL JEM 2010F (Japan) field emission transmission electron microscope with an acceleration voltage of 200 kV. The Seebeck coefficient and resistivity were measured using a ZEM-3 Seebeck coefficient/electric resistance measuring system (Ulvac-Riko, Inc., Japan) with an uncertainty of about 10%. The total thermal conductivities were calculated with the equation κtot = λCpd, where λ, Cp, and d are thermal diffusion coefficient, specific heat capacity, and density, respectively. The thermal diffusion coefficient was evaluated with a laser flash apparatus (TC-9000, Ulvac-Riko, Japan). The specific heat capacities were measured with a DSC STA449 equipment (Netzsch, Germany) as well as referring the calculated value via the Dulong-Petit law. An XS105 density meter (Switzerland) was used to obtain the densities of polished-samples after SPS. The Hall coefficient (RH) was measured using a Hall measurement system (ResiTest 8340DC, Tokyo, Japan), and the Hall carrier concentration (nH) and mobility (μH) were respectively calculated via nH = 1/(eRH) and μH = RH/ρ, in which ρ is the electrical resistivity.

the MnTe NWs were evenly dispersed in the HMS matrix, the κL values at temperatures higher than 525 K were comparable to the κL min of pure HMS, which indicated the incorporation of well-dispersed MnTe nanoinclusions to be an effective strategy to reduce the κL. Although the MnTe nanoinclusions were displayed as nanoparticles both in SEM and TEM (Figure 5A,B,E), their morphologies were recorded on a two-dimensional surface or in a quasi-two-dimentional slice. In fact, as indicated by the 3D simulated structure (Figure 6G), similar nanowire morphologies were still retained after SPS. Also, by taking the intermediary grain boundary into account (Figure 5F), the final MnTe nanoinclusions existed as core/shell NWs in the HMS (shown in Figure 6G), where yellow, red, and blue solid represent the MnTe nanowire, intermediary grain boundary, and HMS matrix, respectively. With this special nano/bulk structure, multiple phonon scattering happened both on the intermediary grain boundary as well as the MnTe nanowire, which finally resulted in a further low κL. Ultimately, the ZT of MnSi1.746Te0.03 sample that using Te NWs as raw material and using the mixture of n-hexane and ethanol as solvent further increased to 0.70 above 773 K (Figure 6H), which is one of the highest values reported in current studies.14,17,18,21−26,29,30,34−37 To further increase the TE performance of HMS, some Te quantum dots, Te nanotubes, or some other nanostructures should be tested.

■

CONCLUSION In this study, a significant enhancement of the ZT value (∼71% increase from 0.41 to 0.70 at 773 and 823 K) was achieved in a novel nano/bulk composite structure by incorporating Te NWs derived MnTe nanophases into HMS. The techniques for preparing pure HMS and synthesizing uniform Te NWs as well as the corresponding surface modification were discussed systematically and in depth. Pure HMS with a relatively high ZT (0.41) was obtained by adjusting the ratio of Mn/Si to 1:1.80 via wet ball milling combined with spark plasma sintering. The ZT increased to 0.60 due to the remarkably reduced κtot, which resulted from porous MnTe inclusions in the MnSi1.746Te0.03 composites. Te NWs with a length/ diameter ratio above 220:1 were successfully synthesized with a facile, efficient, and relatively green solution phase method. Well-dispersed mixtures consisting of Te NWs, Si, and Mn powder were obtained after surface modification on Te NWs. A special MnTe/HMS nano/bulk structure with ∼20 nm intermediary grain boundary capped on the surface of NWs was prepared. κL was further reduced to the theoretical κL min due to the intensified phonon scattering effect. The nano/bulk structure as well as the preparation technologies provided a glimpse of the structure−property relationships of a broad class of materials.

■

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b02270. Additional characterization data, thermoelectric properties of HMS with Si or MnSi phase, electronic probe microanalyzer analysis of MnTe/HMS composite, XRD pattern of MnTe/HMS composite sintered at low temperature, carrier concentration and mobility of pure HMS and MnTe/HMS composite, Seebeck coefficient of pure MnTe, SEM images of MnTe/HMS nano/bulk structures, densities of pure HMS and MnTe/HMS composites (PDF)

EXPERIMENTAL SECTION

Preparation of MnTe/HMS Composites. (MnSi1.80)1−x(MnTe)x composition of the constituent elements including manganese (Mn, ∼400 mesh, 99.9%), silicon (Si, ∼200 mesh, 99.99%), and tellurium (Te, ∼400 mesh, 99.99%) powders were homogeneously mixed via planetary ball milling for 2 h with a rotational speed of 350 rpm in a 250 mL zirconia jar with tungsten carbide balls. Prior to ball milling, ∼40 mL of n-hexane (99.99%) was added to the sealed jar to prevent oxidation and decreased particle sizes. The relatively low ball-topowder weight ratio (approximately 10:1) and a short duration of 2 h divided into four rotating periods in reverse directions were utilized to avoid the formation of an impure MnSi phase. Subsequently, the muddy mixture was adequately dried to powders in a vacuum oven at 7386


Article


■

Thermoelectric Properties of Higher Manganese Silicides. J. Appl. Phys. 2013, 114, 173705. (15) Wu, H. J.; Zhao, L.-D.; Zheng, F. S.; Pei, Y. L.; Tong, X.; Kanatzidis, M. G.; He, J. Q.; Wu, D. Broad Temperature Plateau for Thermoelectric Figure of Merit ZT > 2 in Phase-Separated PbTe0.7S0.3. Nat. Commun. 2014, 5, 4515. (16) 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. (17) Nhi Truong, D. Y.; Kleinke, H.; Gascoin, F. Preparation of Pure Higher Manganese Silicides through Wet Ball Milling and Reactive Sintering with Enhanced Thermoelectric Properties. Intermetallics 2015, 66, 127−132. (18) Chen, X.; Girard, S. N.; Meng, F.; Lara-Curzio, E.; Jin, S.; Goodenough, J. B.; Zhou, J. S.; Shi, L. Approaching the Minimum Thermal Conductivity in Rhenium-Substituted Higher Manganese Silicides. Adv. Energy Mater. 2014, 4, 1400452. (19) Chen, X.; Weathers, A.; Carrete, J.; Mukhopadhyay, S.; Delaire, O.; Stewart, D. A.; Mingo, N.; Girard, S. N.; Ma, J.; Abernathy, D. L.; Yan, J. Q.; Sheshka, R.; Sellan, D. P.; Meng, F.; Jin, S.; Zhou, J. S.; Shi, L. Twisting Phonons in Complex Crystals with Quasi-One-Dimensional Substructures. Nat. Commun. 2015, 6, 6723. (20) Yamamoto, A.; Ghodke, S.; Miyazaki, H.; Inukai, M.; Nishino, Y.; Matsunami, M.; Takeuchi, T. Thermoelectric Properties of Supersaturated Re Solid Solution of Higher Manganese Silicides. Jpn. J. Appl. Phys. 2016, 55, 020301. (21) Itoh, T.; Yamada, M. Synthesis of Thermoelectric Manganese Silicide by Mechanical Alloying and Pulse Discharge Sintering. J. Electron. Mater. 2009, 38, 925−929. (22) Zamanipour, Z.; Shi, X.; Mozafari, M.; Krasinski, J. S.; Tayebi, L.; Vashaee, D. Synthesis, Characterization, and Thermoelectric Properties of Nanostructured Bulk P-Type MnSi1.73, MnSi1.75, and MnSi1.77. Ceram. Int. 2013, 39, 2353−2358. (23) Shi, X.; Shi, X.; Li, Y.; He, Y.; Chen, L.; Li, Q. Enhanced Power Factor of Higher Manganese Silicide via Melt Spin Synthesis Method. J. Appl. Phys. 2014, 116, 245104. (24) Zhou, A. J.; Zhu, T. J.; Zhao, X. B.; Yang, S. H.; Dasgupta, T.; Stiewe, C.; Hassdorf, R.; Mueller, E. Improved Thermoelectric Performance of Higher Manganese Silicides with Ge Additions. J. Electron. Mater. 2010, 39, 2002−2007. (25) Luo, W.; Li, H.; Fu, F.; Hao, W.; Tang, T. Improved Thermoelectric Properties of Al-Doped Higher Manganese Silicide Prepared by a Rapid Solidification Method. J. Electron. Mater. 2011, 40, 1233−1237. (26) Ponnambalam, V.; Morelli, D. T.; Bhattacharya, S.; Tritt, T. M. The Role of Simultaneous Substitution of Cr and Ru on the Thermoelectric Properties of Defect Manganese Silicides MnSiδ (1.73 < δ < 1.75). J. Alloys Compd. 2013, 580, 598−603. (27) Muthiah, S.; Singh, R. C.; Pathak, B. D.; Dhar, A. Facile Synthesis of Higher Manganese Silicide Employing Spark Plasma Assisted Reaction Sintering with Enhanced Thermoelectric Performance. Scr. Mater. 2016, 119, 60−64. (28) Miyazaki, Y.; Hamada, H.; Hayashi, K.; Yubuta, K. Crystal Structure and Thermoelectric Properties of Lightly VanadiumSubstituted Higher Manganese Silicides (Mn1‑xVx)Siγ. J. Electron. Mater. 2017, 46, 2705−2710. (29) Luo, W.; Li, H.; Yan, Y.; Lin, Z.; Tang, X.; Zhang, Q.; Uher, C. Rapid Synthesis of High Thermoelectric Performance Higher Manganese Silicide with in-situ Formed Nano-Phase of MnSi. Intermetallics 2011, 19, 404−408. (30) Girard, S. N.; Chen, X.; Meng, F.; Pokhrel, A.; Zhou, J.; Shi, L.; Jin, S. Thermoelectric Properties of Undoped High Purity Higher Manganese Silicides Grown by Chemical Vapor Transport. Chem. Mater. 2014, 26, 5097−5104. (31) Bernard-Granger, G.; Soulier, M.; Ihou-Mouko, H. I.; Navone, C.; Boidot, M.; Leforestier, J.; Simon, J. Microstructure Investigations and Thermoelectrical Properties of a P-Type Polycrystalline Higher

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Jing-Feng Li: 0000-0002-0185-0512 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS This work was financially supported by the Advanced Material Engineering Division of Toyota Motor Corporation, the National Natural Science Foundation (No. 11474176), the China Postdoctoral Science Foundation (Grant No. 2016M601015), and the National Basic Research Program of China (Grant No. 2013CB632503).

■

REFERENCES

(1) Bell, L. E. Cooling, Heating, Generating Power, and Recovering Waste Heat with Thermoelectric Systems. Science 2008, 321, 1457− 1461. (2) Snyder, G. J.; Toberer, E. S. Complex Thermoelectric Materials. Nat. Mater. 2008, 7, 105−114. (3) Pei, Y. Z.; Shi, X. Y.; LaLonde, A.; Wang, H.; Chen, L. D.; Snyder, G. J. Convergence of Electronic Bands for High Performance Bulk Thermoelectrics. Nature 2011, 473, 66−69. (4) Tan, G. J.; Hao, S. Q.; Zhao, J.; Wolverton, C.; Kanatzidis, M. G. High Thermoelectric Performance in Electron-Doped AgBi3S5 with Ultralow Thermal Conductivity. J. Am. Chem. Soc. 2017, 139, 6467− 6473. (5) Disalvo, F. J. Thermoelectric Cooling and Power Generation. Science 1999, 285, 703−706. (6) Venkatasubramanian, R.; Siivola, E.; Colpitts, T.; O’Quinn, B. Thin-Film Thermoelectric Devices with High Room-Temperature Figures of Merit. Nature 2001, 413, 597−602. (7) 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. (8) Poudel, B.; Hao, Q.; Ma, Y.; Lan, Y. C.; Minnich, A.; Yu, B.; Yan, X.; Wang, D. Z.; Muto, A.; Vashaee, D.; Chen, X. Y.; Liu, J. M.; Dresselhaus, M. S.; Chen, G.; Ren, Z. F. High-Thermoelectric Performance of Nanostructured Bismuth Antimony Telluride Bulk Alloys. Science 2008, 320, 634−638. (9) 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. (10) Biswas, K.; He, J. Q.; 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. (11) Zhao, L. D.; Lo, S. H.; Zhang, Y. S.; Sun, H.; Tan, G. J.; 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. (12) Zhao, L. D.; Tan, G. J.; Hao, S. Q.; He, J. Q.; Pei, Y. L.; Chi, H.; Wang, H.; Gong, S. K.; Xu, H. B.; Dravid, V. P.; Uher, C.; Snyder, G. J.; Wolverton, C.; Kanatzidis, M. G. Ultrahigh Power Factor and Thermoelectric Performance in Hole-Doped Single-Crystal SnSe. Science 2016, 351, 141−144. (13) Migas, D. B.; Shaposhnikov, V. L.; Filonov, A. B.; Borisenko, V. E.; Dorozhkin, N. N. Ab Initio Study of the Band Structures of Different Phases of Higher Manganese Silicides. Phys. Rev. B: Condens. Matter Mater. Phys. 2008, 77, 075205. (14) Chen, X.; Weathers, A.; Salta, D.; Zhang, L.; Zhou, J.; Goodenough, J. B.; Shi, L. Effects of (Al, Ge) Double Doping on the 7387


Article

Chemistry of Materials Manganese Silicide Material Sintered from a Gas-Phase Atomized Powder. J. Alloys Compd. 2015, 618, 403−412. (32) Okamoto, N. L.; Koyama, T.; Kishida, K.; Tanaka, K.; Inui, H. Crystal Structure and Thermoelectric Properties of Chimney-Ladder Compounds in the Ru2Si3-Mn4Si7 Pseudobinary System. Acta Mater. 2009, 57, 5036−5045. (33) Zhou, A. J.; Zhao, X. B.; Zhu, T. J.; Cao, Y. Q.; Stiewe, C.; Hassdorf, R.; Mueller, E. Composites of Higher Manganese Silicides and Nanostructured Secondary Phases and Their Thermoelectric Properties. J. Electron. Mater. 2009, 38, 1072−1077. (34) Zhou, A. J.; Zhao, X. B.; Zhu, T. J.; Yang, S. H.; Dasgupta, T.; Stiewe, C.; Hassdorf, R.; Mueller, E. Microstructure and Thermoelectric Properties of SiGe-Added Higher Manganese Silicides. Mater. Chem. Phys. 2010, 124, 1001−1005. (35) Sadia, Y.; Gelbstein, Y. Silicon-Rich Higher Manganese Silicides for Thermoelectric Applications. J. Electron. Mater. 2012, 41, 1504− 1508. (36) Sadia, Y.; Elegrably, M.; Ben-Nun, O.; Marciano, Y.; Gelbstein, Y. Submicron Features in Higher Manganese Silicide. J. Nanomater. 2013, 2013, 701268. (37) Sadia, Y.; Madar, N.; Kaler, I.; Gelbstein, Y. Thermoelectric Properties of the Quasi-Binary MnSi1.73-FeSi2 System. J. Electron. Mater. 2015, 44, 1637−1643. (38) Chen, X.; Zhou, J.; Goodenough, J. B.; Shi, L. Enhanced Thermoelectric Power Factor of Re-Substituted Higher Manganese Silicides with Small Islands of MnSi Secondary Phase. J. Mater. Chem. C 2015, 3, 10500−10508. (39) Miyazaki, Y.; Igarashi, D.; Hayashi, K.; Kajitani, T.; Yubuta, K. Modulated Crystal Structure of Chimney-Ladder Higher Manganese Silicides MnSiγ (γ ∼ 1.74). Phys. Rev. B: Condens. Matter Mater. Phys. 2008, 78, 214104. (40) Koumoto, K.; Mori, T. Thermoelectric Nanomaterials: Materials Design and Applications; Springer: Berlin, 2013. (41) Ye, H. Q.; Amelinckx, S. High-Resolution Electron Microscopic Study of Manganese Silicides MnSi2‑x. J. Solid State Chem. 1986, 61, 8− 39. (42) Higgins, J. M.; Schmitt, A. L.; Guzei, I. A.; Jin, S. Higher Manganese Silicide Nanowires of Nowotny Chimney Ladder Phase. J. Am. Chem. Soc. 2008, 130, 16086−16094. (43) Sadia, Y.; Dinnerman, L.; Gelbstein, Y. Mechanical Alloying and Spark Plasma Sintering of Higher Manganese Silicides for Thermoelectric Applications. J. Electron. Mater. 2013, 42, 1926−1931. (44) Saleemi, M.; Famengo, A.; Fiameni, S.; Boldrini, S.; Battiston, S.; Johnsson, M.; Muhammed, M.; Toprak, M. S. Thermoelectric Performance of Higher Manganese Silicide Nanocomposites. J. Alloys Compd. 2015, 619, 31−37. (45) An, T. H.; Choi, S. M.; Seo, W. S.; Park, C.; Kim, I. H.; Kim, S. U. The Effect of Microstructure on the Thermoelectric Properties of Polycrystalline Higher Manganese Silicides. Jpn. J. Appl. Phys. 2013, 52, 10MC11. (46) Yamada, T.; Miyazaki, Y.; Yamane, H. Preparation of Higher Manganese Silicide (HMS) Bulk and Fe-Containing HMS Bulk Using a Na-Si Melt and their Thermoelectrical Properties. Thin Solid Films 2011, 519, 8524−8527. (47) Yoshikura, M.; Itoh, T. Thermoelectric Properties of Higher Manganese Silicide Compounds Synthesized by MG-PDS Method. Funtai oyobi Funmatsu Yakin 2010, 57, 242−246. (48) Sadia, Y.; Aminov, Z.; Mogilyansky, D.; Gelbstein, Y. Texture Anisotropy of Higher Manganese Silicide Following Arc-Melting and Hot-Pressing. Intermetallics 2016, 68, 71−77. (49) Chen, X.; Shi, L.; Zhou, J.; Goodenough, J. B. Effects of Ball Milling on Microstructures and Thermoelectric Properties of Higher Manganese Silicides. J. Alloys Compd. 2015, 641, 30−36. (50) Shin, D. K.; Ur, S. C.; Jang, K. W.; Kim, I. H. Solid-State Synthesis and Thermoelectric Properties of Cr-Doped MnSi1.75‑δ. J. Electron. Mater. 2014, 43, 2104−2018. (51) Truong, D. Y. N.; Kleinke, H.; Gascoin, F. Thermoelectric Properties of Higher Manganese Silicide/Multi-Walled Carbon Nanotube Composites. Dalton Trans. 2014, 43, 15092−15097.

(52) Shin, D. K.; You, S. W.; Kim, I. H. Solid-State Synthesis and Thermoelectric Properties of Al-Doped MnSi1.73. J. Korean Phys. Soc. 2014, 64, 1412−1415. (53) Nhi Truong, D. Y.; Berthebaud, D.; Gascoin, F.; Kleinke, H. Molybdenum, Tungsten, and Aluminium Substitution for Enhancement of the Thermoelectric Performance of Higher Manganese Silicides. J. Electron. Mater. 2015, 44, 3603−3611. (54) Wu, G. R.; Nagatomo, K.; Sasaki, M.; Nagasaki, F.; Sato, H.; Taniguchi, M.; Gao, W. X. Photo-Induced Transient Thermoelectric Effect in MnTe. Solid State Commun. 2001, 118, 425−429. (55) Zhang, L. L.; Wang, W.; Ren, B.; Yan, Y. Thermoelectric Performance and High-Temperature Creep Behavior of GeTe-Based Thermoelectric Materials. J. Electron. Mater. 2011, 40, 1057−1061. (56) Kim, B.; Kim, I.; Min, B. K.; Oh, M.; Park, S.; Lee, H. Thermoelectric Properties of Non-Stoichiometric MnTe Compounds. Electron. Mater. Lett. 2013, 9, 477−480. (57) Xie, W.; Populoh, S.; Galazka, K.; Xiao, X.; Sagarna, L.; Liu, Y.; Trottmann, M.; He, J.; Weidenkaff, A. Thermoelectric Study of Crossroads Material MnTe via Sulfur Doping. J. Appl. Phys. 2014, 115, 103707. (58) Wu, H. J.; Chang, C.; Feng, D.; Xiao, Y.; Zhang, X.; Pei, Y.; Zheng, L.; Wu, D.; Gong, S.; Chen, Y.; He, J.; Kanatzidis, M. G.; Zhao, L. D. Synergistically Optimized Electrical and Thermal Transport Properties in SnTe via Alloying High-Solubility MnTe. Energy Environ. Sci. 2015, 8, 3298−3312. (59) Tan, G. J.; Shi, F. Y.; Hao, S. Q.; Chi, H.; Bailey, T. P.; Zhao, L. D.; Uher, C.; Wolverton, C.; Dravid, V. P.; Kanatzidis, M. G. Valence Band Modification and High Thermoelectric Performance in SnTe Heavily Alloyed with MnTe. J. Am. Chem. Soc. 2015, 137, 11507− 11516. (60) Abrikosov, N. K.; Dyuldina, K. A.; Zhdanova, V. V. Temperature-Composition Phase Diagrams of the Systems-Lead Selenide-Mercury Selenide and Lead Telluride-Mercury. Inorg. Mater. 1968, 4, 1638−1642. (61) Vanyarkho, V. G.; Zlomanov, V. P.; Novoselova, A. V. Lead Telluride-Manganese Telluride System. Inorg. Mater. 1970, 6, 1534− 1539. (62) Mamontov, M. N. Phase Relationships in the CdTe-MnTe-Te System at 643 K. Inorg. Mater. 1996, 32, 716−720. (63) Wei, T.-R.; Tan, G. J.; Zhang, X. M.; Wu, C.-F.; Li, J.-F.; Dravid, V. P.; Snyder, G. J.; Kanatzidis, M. G. Distinct Impact of Alkali-Ion Doping on Electrical Transport Properties of Thermoelectric P-Type Polycrystalline SnSe. J. Am. Chem. Soc. 2016, 138, 8875−8882. (64) Dresselhaus, M. S.; Chen, G.; Tang, M. Y.; Yang, R.; Lee, H.; Wang, D. Z.; Ren, Z. F.; Fleurial, J. P.; Gogna, P. New Directions for Low-Dimensional Thermoelectric Materials. Adv. Mater. 2007, 19, 1043−1053. (65) Faleev, S. V.; Léonard, F. Theory of Enhancement of Thermoelectric Properties of Materials with Nanoinclusions. Phys. Rev. B: Condens. Matter Mater. Phys. 2008, 77, 214304. (66) Martin, J.; Wang, L.; Chen, L. D.; Nolas, G. S. Enhanced Seebeck Coefficient through Energy-Barrier Scattering in PbTe Nanocomposites. Phys. Rev. B: Condens. Matter Mater. Phys. 2009, 79, 115311. (67) Yang, X. H.; Qin, X. Y. Giant Scattering Parameter and Enhanced Thermoelectric Properties Originating from Synergetic Scattering of Electrons in Semiconductors with Metal Nanoinclusions. Appl. Phys. Lett. 2010, 97, 192101. (68) Li, J. H.; Tan, Q.; Li, J.-F.; Liu, D. W.; Li, F.; Li, Z. Y.; Zou, M. M.; Wang, K. BiSbTe-Based Nanocomposites with High ZT: the Effect of SiC Nanodispersion on Thermoelectric Properties. Adv. Funct. Mater. 2013, 23, 4317−4323. (69) Li, Z.; Xiao, C.; Zhu, H.; Xie, Y. Defect Chemistry for Thermoelectric Materials. J. Am. Chem. Soc. 2016, 138, 14810−14819. (70) Tang, J. Y.; Wang, H.-T.; Lee, D. H.; Fardy, M.; Huo, Z. Y.; Russell, T. P.; Yang, P. D. Holey Silicon as an Efficient Thermoelectric Material. Nano Lett. 2010, 10, 4279−4283. (71) Lee, J.; Lim, J.; Yang, P. D. Ballistic Phonon Transport in Holey Silicon. Nano Lett. 2015, 15, 3273−3279. 7388


Article

Chemistry of Materials (72) Vineis, C. J.; Shakouri, A.; Majumdar, A.; Kanatzidis, M. G. Nanostructured Thermoelectrics: Big Efficiency Gains from Small Features. Adv. Mater. 2010, 22, 3970−3980. (73) Li, Z. L.; Zheng, S. Q.; Zhang, Y. Z.; Teng, R. Y.; Huang, T.; Chen, C. F.; Lu, G. W. Controlled Synthesis of Tellurium Nanowires and Nanotubes via a Facile, Efficient, and Relatively Green Solution Phase Method. J. Mater. Chem. A 2013, 1, 15046−15052. (74) Li, Z. L.; Zheng, S. Q.; Zhang, Y. Z.; Chen, H.; Huang, T.; Lu, G. W. High-Yield Synthesis, Controllable Evolution, and Thermoelectric Properties of Te/Bi2Te3 Heterostructure Nanostrings. J. Electron. Mater. 2015, 44, 2061−2067. (75) Cahill, D. G.; Watson, S. K.; Pohl, R. O. Lower Limit to the Thermal Conductivity of Disordered Crystals. Phys. Rev. B: Condens. Matter Mater. Phys. 1992, 46, 6131.

7389


Significant Enhancement of the Thermoelectric Performance of Higher

Recommend Documents