High Thermoelectric Figure of Merit Achieved in Cu2S1âxTex Alloys

Research Article Cite This: ACS Appl. Mater. Interfaces 2018, 10, 32201−32211

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High Thermoelectric Figure of Merit Achieved in Cu2S1−xTex Alloys Synthesized by Mechanical Alloying and Spark Plasma Sintering Yao Yao,† Bo-Ping Zhang,*,† Jun Pei,† Qiang Sun,‡ Ge Nie,§ Wen-Zhen Zhang,† Zhen-Tao Zhuo,† and Wei Zhou† †

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The Beijing Municipal Key Laboratory of New Energy Materials and Technologies, School of Materials Science and Engineering, University of Science and Technology Beijing, 100083 Beijing, China ‡ Department of Materials Science and Engineering, COE, Peking University, Beijing 100871, China § ENN Group, Langfang City, Hebei Province 065001, China S Supporting Information *

ABSTRACT: Chalcogenides have been considered as promising thermoelectric materials because of their low cost, nontoxicity, and environmental benignity. In this work, we synthesized a series of Cu2S1−xTex (0 ≤ x ≤ 1) alloys by a facile, rapid method of mechanical alloying combined with spark plasma sintering process. The Cu2S1−xTex system provides an excellent vision of the competition between pure phase and phase transformation, entropy-driven solid solution, and enthalpydriven phase separation. When the Te concentration increases, the Cu2S1−xTex system changed from the pure monoclinic Cu2S at x = 0 to monoclinic Cu2S1−xTex solid solution at 0.02 ≤ x ≤ 0.06 and then transforms to hexagonal Cu2S1−xTex solid solution at 0.08 ≤ x ≤ 0.1. The phase separation of hexagonal Cu2Te in the hexagonal Cu2S matrix occurs at 0.3 ≤ x ≤ 0.7 and finally forms the hexagonal Cu2Te at x = 1. Owing to the changed band structure and the coexisted Cu2S and Cu2Te phases, greatly enhanced power factor was achieved in all Cu2S1−xTex (0 < x < 1) alloys. Meanwhile, the point defect introduced by the substitution of Te/S atoms strengthened the phonon scattering, resulting in a lowered lattice thermal conductivity in most of these solid solutions. As a consequence, Cu2S0.94Te0.06 exhibits a maximum ZT value of 1.18 at 723 K, which is about 3.7 and 14.8 times as compared to the values of pristine Cu2S (0.32) and Cu2Te (0.08), respectively. KEYWORDS: thermoelectric, Cu2S, Cu2Te, mechanical alloying, spark plasma sintering cost, nontoxicity, and environmental benignity.16 Copper sulfide and copper telluride are both important semiconductor copper chalcogenides, which have complex crystal structures accompanied by several phase transitions. Cu2S shows three kinds of crystal structures with increasing temperature: lowtemperature monoclinic structure (below 370 K), midtemperature hexagonal structure (370−700 K), and hightemperature cubic structure (above 700 K).17 Compared to Cu2S, the Cu2Te compound exhibits more complicated temperature-dependent crystal structures including hexagonal α phase (below 454 K), hexagonal β phase (454−597 K), orthorhombic γ phase (597−627 K), hexagonal δ structure (627−756 K), and cubic ε phase (above 756 K).18 Liu et al.19 reported that the occurrence of phase transitions could cause critical scattering for both charge carriers and phonons, which is favorable for improving ZT. In recent years, the TE properties of Cu2S and Cu2Te have been extensively studied, such as adjusting the copper deficiency in Cu2S to optimize the

1. INTRODUCTION Thermoelectric (TE) materials can convert thermal energy to electrical power directly and vice versa, which are expected to play an important role in power generation and cooling.1−3 Because of the advantages such as no moving parts, no noise, and exceptional reliability, TE materials show great potential for solving the global energy crisis and environmental pollution.4 The energy conversion efficiency of TE materials is determined by the dimensionless figure of merit ZT = α2σT/ (κe + κl), where α is the Seebeck coefficient, σ is the electrical conductivity, T is the absolute temperature, κe is the electronic thermal conductivity, and κl is the lattice thermal conductivity.5 Generally, there are two strategies to increase the ZT of TE materials: increasing the power factor (α2σ)6 using band convergence,7 quantum confinement effect,8 resonance level,9 and carrier concentration regulation10 by doping and reducing the thermal conductivity (κ) via point defect scattering by solid solution,11 nanostructuring,12 mesoscale grain boundaries,13 hierarchical architecture engineering,14 and defect engineering.15 Chalcogenide materials have been widely investigated as promising TE materials in recent years because of their low © 2018 American Chemical Society

Received: July 6, 2018 Accepted: September 4, 2018 Published: September 4, 2018 32201

DOI: 10.1021/acsami.8b11300 ACS Appl. Mater. Interfaces 2018, 10, 32201−32211

Research Article

ACS Applied Materials & Interfaces

Figure 1. XRD patterns of bulk samples within 2θ range of 20°−50° for Cu2S1−xTex (a) x = 0, 0.02, 0.04, 0.06, and 0.08 and (b) x = 0.1, 0.3, 0.5, 0.7, and 1. (c) 2θ range of 23°−27° at x = 0.1, 0.3, 0.5, 0.7, and 1.

hole concentration,20−22 doping element to modulate band structure or form composite,23,24 fabricating mosaic structure with a mixture of elements with vastly mismatched half-S and half-Te to form Cu2S0.5Te0.5 solid solution,25 codoping Ni and Se elements in Cu2Te to increase the carrier concentration and reduce the thermal conductivity,26 and so forth. However, the TE properties of Te-doped Cu2S alloys with different S/Te atomic ratios have rarely been reported so far. Considering the low σ and κ values of Cu2S and high σ value of Cu2Te, the Cu2S1−xTex alloys are expected to achieve an enhanced TE performance. In addition, the point defects and mass fluctuations introduced by S/Te alloying would effectively scatter the phonons, resulting in a reduction of κl. Hence, it would be interesting to investigate the phase structure and TE properties of ternary Cu2S1−xTex alloys with different S/Te atomic ratios. Besides the improvement strategies, the preparation technology is also very important for the enhancement and practical applications of TE materials. It is well known that many preparation processes of TE materials are timeconsuming and energy-intensive, which involves melting, crushing, and annealing for several days.27 Mechanical alloying (MA) combined with spark plasma sintering (SPS) technique has a big advantage in shortening the preparation process compared with those complex methods, which is beneficial for practical applications. The MA process provides high energy to the powder particles through the inertia of the grinding balls, which includes cold welding, fracturing, rewelding, and repulverization. The SPS then accomplishes the sintering process within several minutes, retaining the structures inherited from the MA-derived powders.28 The convenient and rapid MA & SPS method has been applied for many highperformance TE materials such as Bi2Te3-based alloys,29 PbTebased alloys,30 selenides,31,32 and sulfides,33,34 which provides us a reliable method to synthesize Cu2S1−xTex compounds. Besides that, the high-energy ball milling and quick hot pressing method is a more popular and effective fabrication method to synthesize materials such as BiSbTe,35 AgSbSe2,36 and GeTe,37 especially for some compounds that are really difficult to prepare using the traditional methods such as skutterudite38 and α-MgAgSb.39

In this work, we synthesized a series of Cu2S1−xTex (x = 0, 0.02, 0.04, 0.06, 0.08, 0.1, 0.3, 0.5, 0.7, and 1) compounds by a facile, rapid method of MA combined with SPS technique. The phase composition, crystal structure, microstructure, and TE properties have been systematically studied with an emphasis on the effect of x, and the relevant mechanisms are discussed. Both enhanced power factor and decreased thermal conductivity result in an improved ZT of 1.18 for Cu2S0.94Te0.06 at 723 K, which is about 3.7 times that (0.32) of pristine Cu2S and 14.8 times that (0.08) of pristine Cu2Te.

2. EXPERIMENTAL SECTION Polycrystalline samples were fabricated via MA combined with SPS. Commercial elemental powders of Cu (99.5%), S (99.8%), and Te (99.9%) were weighed according to the stoichiometric proportions of CuS1−xTex (x = 0, 0.02, 0.04, 0.06, 0.08, 0.1, 0.3, 0.5, 0.7, and 1) and milled on a planetary ball mill (QM-3SP2, Nanjing University, China) at 425 rpm for 2 h filled with a mixed atmosphere of 95 vol % argon and 5 vol % hydrogen. Stainless steel vessel and balls are used, and the weight ratio of ball to powders was kept at 20:1. The MA-derived powders were loaded into a Φ 20 mm graphite mold and then sintered at 773 K for 5 min under axial pressure of 40 MPa in vacuum using a SPS system (SPS-211Lx, Fuji Electronic Industrial, Japan), yielding a disk-shaped bulk ∼3 mm in height and ∼20 mm in diameter. The Phase structure was investigated by X-ray diffraction (XRD, Rigaku 2500, Japan) with Cu Kα radiation. The microstructure of the bulks was observed using a field emission scanning electron microscopy (FESEM, SUPRA 55, Germany). The transmission electron microscopy (TEM) investigation was conducted using a JEOL 2100F microscope operated at 200 kV. Optical absorption measurement was conducted on bulk samples by using the UV−vis− NIR spectrum (Cary 5000, Varian, America) with an integrating sphere at room temperature. The optical band gaps (Eg’s) were estimated by extrapolating (αhν)n to 0 as a function of hν, where α, h, and ν are the absorption coefficient, Plank constant, and light frequency, respectively. n is equal to 2 for direct gaps and 1/2 for indirect gaps, which n = 2 in our work owing to the direct gap of Cu2S.40 Electrical conductivity (σ) and Seebeck coefficient (α) as a function of temperature were conducted on a bar-shaped specimen from 323 to 773 K in a static helium atmosphere using a LSR-3 electric resistance/Seebeck coefficient measuring system (LSR-3, LINSEIS, Germany). Carrier concentration (n) and mobility (μ) were measured under a reversible magnetic field using a Hall measurement system (ResiTest 8400, Toyo, Japan) at room temperature. Thermal diffusivity (D) and heat capacity (Cp) were measured using the laser 32202


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Figure 2. Lattice parameters (a), density and relative density (b) for Cu2S1−xTex (x = 0, 0.02, 0.04, 0.06, 0.08, 0.1, 0.3, 0.5, 0.7, and 1) bulks, and crystal structures for M-Cu2S, H-Cu2S (c). flash method (NETZSCH, LFA457, Germany) at temperatures from 323 to 773 K. The thermal conductivity (κ) was calculated using the equation κ = DdCp, with mass density (d) of the sample measured by the Archimedes method.

patterns at 23°−27°, as shown in Figure 1c. It indicates that the solubility of Cu2Te into Cu2S is below 30% in Cu2S1−xTex prepared by applying the MA & SPS process. The diffraction peak intensity of H-Cu2Te becomes stronger as increasing x, suggesting an increased phase ratio of H-Cu2Te to H-Cu2S with x in Cu2S1−xTex at 0.3 ≤ x ≤ 0.7. A rough fraction of HCu2Te to H-Cu2S was obtained through comparing the intensity of the strongest diffraction peak of each phase, and the percentage of H-Cu2Te/(H-Cu2S + H-Cu2Te) increases from 8.9 to 18.1% and 32.5% when x was increased from 0.3 to 0.5 and 0.7. At x = 1, the diffraction peaks can be indexed as hexagonal Cu2Te with the space group P3m1(156) (H-Cu2Te, PDF#49-1411, a = b = 8.357 Å, and c = 21.63 Å) and hexagonal Cu2−xTe with the space group P3m1(156) (HCu2−xTe, PDF#10-0421, a = b = 8.342 Å, and c = 21.69 Å). The room-temperature crystal structure has also been in controversy for decades. In a previous study, it was reported that the SPS sample and directly annealed samples all consisted of more than one crystal structure.18,42 Apart from the phase transition and varied phase fraction, obvious diffraction peak shift with increasing x is also observed, which implies the occurrence of crystal cell volume variation in Cu2S1−xTex. Considering the larger ionic radius of Te2− (2.21 Å) than that of S2− (1.84 Å), such a peak shift proves an evidence of the mutual substitution between Te and S. The shift of the XRD patterns and the correspondingly changed lattice parameters also implies the mass and strain fluctuation after Te alloying. On the basis of the XRD patterns, the lattice parameters (a, b, and c) of Cu2S1−xTex bulks are obtained as shown in Figure 2a, in which the phase compositions change with x: pure MCu2S at 0 ≤ x ≤ 0.06, pure H-Cu2S at 0.08 ≤ x ≤ 0.10, coexistent H-Cu2S and H-Cu2Te at 0.3 ≤ x ≤ 0.7, and HCu2Te at x = 1. The corresponding crystal structures of MCu2S and H-Cu2S are shown in Figure 2c. Because of the larger ionic radius of Te2− (2.21 Å) for S2− (1.84 Å), the a, b, and c for both M-Cu2S (0 ≤ x ≤ 0.06) and H-Cu2S (0.08 ≤ x ≤ 0.7) increase with increasing x. At x = 1, Cu2Te consists of two crystal structures, which are very similar. Thus, it is difficult to

3. RESULTS AND DISCUSSION The XRD patterns of Cu2S1−xTex (x = 0, 0.02, 0.04, 0.06, 0.08, 0.1, 0.3, 0.5, 0.7, and 1) powder samples after milling are shown in the Supporting Information (Figure S1). The Cu2S1−xTex powders have similar diffraction peaks at 0 ≤ x ≤ 0.1, which consist of monoclinic Cu2S (○) with the space group P21/c (M-Cu2S, PDF#83-1462, a = 15.246 Å, b = 11.884 Å, and c = 13.494 Å) and tetragonal Cu1.96S (▼) with the space group P43212 (T-Cu1.96S, PDF#29-0578, a = b = 3.996 Å, and c = 11.287 Å). When x was increased to 0.3 ≤ x ≤ 0.7, some peaks of M-Cu2S and T-Cu1.96S disappeared, whereas peaks from hexagonal Cu2−xTe (□) with the space group P3m1(156) (H-Cu2−xTe, PDF#10-0421, a = b = 8.342 Å, and c = 21.69 Å) arise, which stems from the large amount of Te addition. The diffraction peaks at x = 1 are well matched with those of H-Cu2−xTe, indicating the formation of pure Cu2−xTe phase. Figure 1 shows the XRD patterns of Cu2S1−xTex bulk samples by applying SPS. The diffraction peaks of x = 0 bulk were indexed to M-Cu2S (○) without any secondary phases within the detectable limit, indicating that the single M-Cu2S phase can be synthesized at x = 0 by combining MA & SPS technique, which has reported in our previous work.23 After adding Te, M-Cu2S (○) remains for 0.02 ≤ x ≤ 0.06 and becomes a pure hexagonal structure Cu2S (●) with the space group P63/mmc (194) (H-Cu2S, PDF#26-1116, a = b = 4.033 Å, and c = 6.739 Å) when 0.08 ≤ x ≤ 0.10. The phase transition from M-Cu2S (○) to H-Cu2S (●) stems from the more Te incorporated into the Cu2S matrix. Similar phase transformation has been reported in the Cu1.8S1−xSex system, in which the low-temperature H-Cu1.8S phase changes to the high-temperature cubic structure (C-Cu1.8S) in Cu1.8S1−xSex as Se content x ≥ 0.3.41 As x increases to 0.3, phase separation of Cu2Te phase (☆) appears, which is clearer in the enlarged 32203


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Figure 3. SEM images of the fractured surface morphologies for the Cu2S1−xTex bulks, (a) x = 0, (b) x = 0.02, (c) x = 0.04, (d) x = 0.06, (e) x = 0.08, (f) x = 0.1, (g) x = 0.3, (h) x = 0.5, (i) x = 0.7, and (j−l) x = 1.

obtain the accurate lattice parameter for each phase. However, we can also infer the lattice variation from the shift of diffraction peak. The peak position of Cu2Te is 25.255° at x = 0.3 and larger than that (24.581°) of x = 1 sample, implying a shrunken lattice caused by the replacement of S2− for Te2− in Cu2Te. With increasing x, the diffraction peak of Cu2Te shifts to lower angle, which means that the S2− substitution content becomes less. Simply put, S2− substitutes Te2− in Cu2Te and Te2− displaces S2− in Cu2S when 0.3 ≤ x ≤ 0.7, that is, the S/Te bidirectional substitution and can be expressed by eq 1. Cu 2S + Cu 2Te → Cu 2S1 − x Tex + Cu 2Te1 − xSx

(1)

Figure 2b gives the theoretical density, the measured density, and the relative density for all of the Cu2S1−xTex bulks. The theoretical density is 5.78, 5.78, and 7.33 g/cm3 for the pure M-Cu2S, H-Cu2S, and H-Cu2Te,42 respectively, and theoretical density with rough calculation. The gray area in Figure 2b shows the approximate extent of relative density, which is above 86%, revealing the well compactness of all of the compounds. Figure 3 shows the SEM micrographs of the fractured surfaces for the Cu2S1−xTex bulks. The morphology of x = 0 is shown in Figure 3a. After adding Te, more pores appeared at 0.02 ≤ x ≤ 0.06 bulks, whose crystal structure remains MCu2S. Because of the similar melting point between Cu2S (∼1130 °C) and Cu2Te (∼1140 °C),17 the increased porosity can be attributed to the Te volatilization during SPS processing, which caused by the weaker Cu−Te bond than Cu−S bonds. Meanwhile, the grain grows at 0.02 ≤ x ≤ 0.1, owing to the small diffusion activation energy and large diffusion rate of the substitution solid solution.43 With further increasing 0.3 ≤ x ≤ 0.7, a dense structure and refined grain are observed in Figure 3g−i, which can be ascribed to the coexisted H-Cu2S and H-Cu2Te phases. The x = 1 bulk shows a layered feature in Figure 3j−l, which was squashed along the direction perpendicular to the SPS pressure in Figure 4j and show a lamellar morphology parallel to the SPS compressing direction in Figure 3k. The thickness of layer ranges tens to hundreds of nanometers, as shown in the magnified SEM image in Figure 3l. TEM and high-resolution TEM (HRTEM) observations were further conducted on the x = 0.06 bulk, and the images

Figure 4. TEM and HRTEM images of the x = 0.06 bulks, (a) grains, (b) distorted grain boundaries, (c) twin crystals, (d) copious lattice orientations, (e) enlarged HRTEM image of selected region from (d) and FFT diffractogram, and (f) lattice defects.

are given in Figure 4, where the morphology of grains with a size of 50−100 nm (Figure 4a) is different from the large grain more than 2 μm as seen in the SEM of Figure 3d, indicating that many small grains form a big grain or particle. Obvious distortion in the grain boundary appeared between the little grains, as shown in Figure 4b, and a structure of twin crystals is observed in Figure 4c. The HRTEM image in Figure 4d exhibits copious lattice orientations, which means a disorder growth pattern and demonstrates the polycrystalline crystal. The enlarged HRTEM image of the selected region from Figure 4d is shown in Figure 4e, indicating that the lattice spacings of 0.248 and 0.222 nm for M-Cu2S correspond to the (−612) and (−152) planes, respectively, together with the fast 32204


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Figure 5. Plots of (F(R) × hν)2 vs photon energy (hν) (a), schematic diagram of band structure (b), the carrier concentration (n) and mobility (μ) as a function of x (c) for the Cu2S1−xTex bulk samples.

increased n. The Hall experiments exhibit that the n increases gradually with increasing x as shown in Figure 5c, going from 5.3 × 1018 cm−3 for x = 0 to 5.5 × 1021 cm−3 for x = 1. Because Te2− and S2− are equivalent, excess hole should not be introduced by ionic substitution. Hence, we deem that the increased n is attributed to the smaller band gap with Te doping. At x ≥ 0.3, the n in Figure 5c increases sharply, which arises from the appearance of high n phase Cu2Te. The ultrahigh n of Cu2Te is mainly derived from its severe Cu deficiency in the matrix,42 resulting in 3 orders of magnitude larger n than Cu2S. The variation of carrier mobility (μ) is related to its DOS effective mass (m*) as follows. Figure 6 gives the Seebeck coefficient (α) as a function of n to evaluate the m* at room temperature. On the basis of the

Fourier transform (FFT) pattern in the inset. In addition, the areas within the white marking in Figure 4f reveal the vast lattice defects including the dislocations and lattice distortions, which originates from the synthesis pressure and Te doping. The defected microstructure could effectively scatter the phonons and contribute to the decreased thermal conductivity.44 The band gap (Eg) of TE materials is one of the key factors to affect the carrier transport and the TE properties. Zeier et al.45 reported that the Eg of a compound is related to the different energies of the atomic orbitals, that is, the larger difference in electronegativity of cationic and anionic elements, the larger the Eg is. As the smaller electronegativity of Te (2.10) than that of S (2.58), smaller Eg will be obtained in Cu−Te bonds than that of Cu−S bonds. The reported Eg = 1.04 eV for Cu2Te25 and Eg = 1.20 eV for Cu2S46 give a good verification. Hence, in our case, the Eg of Cu2S will become narrow when doping Te. Even though the crystal structure changed to H-Cu2S when x ≥ 0.08 from M-Cu2S, the Eg should also shows a downtrend, owing to the smaller Eg of HCu2S than that of M-Cu2S as reported earlier.22 In addition, the contribution of Te2− 5p electrons could cause an increase in the total density of states (DOS) of Cu2S1−xTex alloys, resulting in a smaller Eg. Zhao et al.47 also reported an increased total DOS in Cu2Se introduced by the substitution of Se2− by Te2−, which gives rise to a narrowed Eg and an incremental electrical conductivity. In order to verify the change of Eg in our work, an optical absorption measurement was conducted on the representative samples at 0 ≤ x ≤ 0.08. The Eg can be obtained by extrapolating the straight line portion of (αhν)2 versus hν as shown in Figure 5a, in which the Eg values are 1.10, 0.98, 0.95, and 0.84 eV at x = 0, 0.04, 0.06, and 0.08, respectively. This proves that the Eg decreases with increasing x, and the change of band structure can be illustrated, as shown in Figure 5b. The decreased Eg is benefit to the carrier transport, which will have an influence on the carrier concentration (n). Besides that, the changed concentration of Cu vacancies and Burstein−Moss shift can also cause the Fermi level to move to the valence band, contributing to an

Figure 6. Relationship between the Seebeck (α) and carrier concentration (nH) at room temperature.

single parabolic band model, m* can be estimated by the Pisarenko relation, which is expressed as48

8π 2kB 2 ji π zy2/3 jj zz m*T (2) 3eh2 k 3n { where kB is the Boltzmann constant, e is the electron charge, and h is the Planck constant. The theoretical data given by earlier research25 were depicted by red and blue lines, which α=

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function of x show an opposite trend compared with the σ, that is, M-Cu2S has higher α and H-Cu2Te has lower α. The thermal conductivity (κ) and ZT are shown in Figure 7b. The κ at x = 0 is 0.44 W m−1 K−1, and it has a little decline at 0.02 ≤ x ≤ 0.1, which is owing to the strengthened phonon point defect scattering. It is noteworthy that the little different κ at 0 ≤ x ≤ 0.1, which implies that the changed crystal structure from M-Cu2S at 0 ≤ x ≤ 0.06 to H-Cu2S at 0.08 ≤ x ≤ 0.1 has little impact on the κ. However, the κ at 0.3 ≤ x ≤ 1 shows an obviously increased trend owing to the appearance of Cu2Te, in which the high electrical thermal transport of Cu2Te contributes much to the κ. On the basis of the measured electrical and thermal transport properties, the ZT at 323 K was calculated and listed in Figure 7b, whose highest ZT value of 0.19 achieved at x = 0.06. The temperature dependence of the TE transport properties for all of the Cu2S1−xTex bulks was also measured and shown in Figure 8, and the data of Cu2S and Cu2Te from

represents 0.5 me and 4.5 me, respectively. It is should be noted here that the scattering mechanism of alloying and ionized impurity may be also contained in this system, even though we assume that the acoustic phonons dominate the scattering of charge carriers. The x = 0 and x = 1 samples in our work show values similar to other reports,25 which is ca. 0.5 me. The m* at 0.02 ≤ x ≤ 0.06 is higher than that at x = 0, suggesting an affected band structure of M-Cu2S for the 0.02 ≤ x ≤ 0.06 bulks. Correspondingly, the μ value in Figure 5c at 0.02 ≤ x ≤ 0.06 decreases with increasing x according to μ ∝ 1/m*. With increasing Te content, the samples at 0.08 ≤ x ≤ 0.7 show even larger m*, which can be ascribed to the transformed crystal structure from M-Cu2S to H-Cu2S. It makes that x = 0.08 and 0.1 samples have lower μ owing to the higher m* for H-Cu2S. However, the μ increases with x at 0.3 ≤ x ≤ 0.7, owing to the influence of the appearance of Cu2Te, which has much larger μ than Cu2S, as shown in Figure 5c. The large μ in Cu2Te is attributed to the weaker chemical bonds for Cu−Te than that of Cu−S because Te is less electronegative than S.42 In order to understand the influences of crystal structure along with the Te content x on the TE performance of Cu2S1−xTex bulks, the TE transport properties for all samples at 323 K are given in Figure 7. The electrical conductivity (σ)

Figure 7. Electrical conductivity σ and Seebeck coefficient α (a), and thermal conductivity κ and ZT (b) as a function of x at 323 K for the Cu2S1−xTex bulk samples. Figure 8. Temperature dependence of electrical conductivity σ (a), Seebeck coefficient α (b), and power factor α2σ (c) for the Cu2S1−xTex bulk samples.

−1

is 5.6 S cm at x = 0 and has a rising trend with increasing x at 0 ≤ x ≤ 0.06 (M-Cu2S). A slight decline on σ with increasing x is observed at 0.08 ≤ x ≤ 0.1, which belongs to the pure HCu2S. This indicates that H-Cu2S has lower σ than M-Cu2S. Further increasing x produces an increased σ to 52.1, 129.4, and 935.6 S cm−1 at x = 0.3, 0.5, and 0.7, respectively, which is ascribed to the appearance of high σ phase Cu2Te. At x = 1, the Cu2Te phase reaches an extremely high σ value even up to 7666.8 S cm−1 owing to its intrinsically large amount of Cu vacancy. Correspondingly, the Seebeck coefficients (α) as a

literatures18,25 were also listed for comparison. The electrical conductivity (σ) in Figure 8a of x = 0 has two turning points at 370 and 700 K, which correspond to the two-phase transition from M-Cu2S to H-Cu2S and to C-Cu2S. The monoclinic phase below 370 K shows an increased σ with an increase of temperature, and the cubic phase above 700 K shows a decreased σ as increasing temperature. These transport 32206


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Figure 9. Temperature dependence of thermal conductivity κ (a), Lorenz number L (b), electronic thermal conductivity κe (c), lattice thermal conductivity κl (d), scattering parameters Γ (e), and κl at 323 K (red dots) and calculated κl (black line) (f) for the Cu2S1−xTex bulks.

also takes place in Cu2S1−xIx and Cu2S1−xSex systems.19,23 On the basis of the σ and α, the power factors (PF = σ2α) for all samples were calculated and are shown in Figure 8c. The PF of most Te-doped Cu2S samples is greatly improved. The maximum PF is 953.9 μW m−1 K−2 at 773 K when x = 0.7, which is 73% higher than that (550.1 μW m−1 K−2) of Cu2Te and increases 5.5 times than that (146.1 μW m−1 K−2) of Cu2S. Figure 9 shows the temperature dependence of thermal transport properties for the Cu2S1−xTex samples. The κ for all of the samples in Figure 9a exhibits a temperature-independent character in the whole measuring temperature range, except a few turning points among 573−773 K at x = 1, which correspond to the phase transitions of Cu2Te at 573, 627, and 756 K.18 The κ at x = 0 ranges 0.42−0.45 W m−1 K−1 at 323− 773 K and becomes lower after doping Te at 0.02 ≤ x ≤ 0.1 and turns to higher when x ≥ 0.3. The κ is composed of the electronic thermal conductivity (κe) and the lattice thermal conductivity (κ l ). The κ e can be calculated by the Wiedemann−Franz law: κe = LσT, where L is the Lorenz number and T is the absolute temperature. The L can be estimated from the following equations assuming a single parabolic band mode49

properties suggest a typical intrinsic semiconducting behavior for M-Cu2S and metallic behavior for C-Cu2S. The σ at x = 0 ranges from 3.2 to 10.8 S cm−1 in the whole measuring temperature. With increasing x, the σ increases gradually at 0.02 ≤ x ≤ 0.06. The increased σ mainly attributes to the changed band structure, which reduces the band gap and facilitates the carrier transport. As increasing x = 0.08 and x = 0.1, σ has a slight decrease, which ascribes to the varied HCu2S from M-Cu2S. Further increasing x produces a sharply increased σ at 0.03 ≤ x ≤ 0.07, owing to the appearance of high σ phase Cu2Te. At x = 1, the Cu2Te phase reaches an ultrahigh σ value up to 7666.8 S cm−1 at 323 K. All samples exhibit positive values of Seebeck coefficients (α) over the entire temperature range as shown in Figure 8b, showing a p-type character whose dominant charge carrier is holes. In contrast with the behavior of σ, the α decreases with increasing x at 0 ≤ x ≤ 0.06 and then has a slight upswing when x = 0.08 and x = 0.1 and finally drops markedly when x ≥ 0.3. The α increases with increasing temperature at 323− 773 K for all samples. It is noteworthy that the abrupt increase in α around 723 K for 0.02 ≤ x ≤ 0.06 samples could be associated with the enhanced critical scattering in the phase transition induced by the moderate substitution of Te for S. The similar extreme variation at phase transition temperature 32207


ACS Applied Materials & Interfaces Ä Å 2Å ij kB yz ÅÅÅ (r + 7/2)Fr + 5/2(ξ) j z L = jj zz ÅÅÅ k e { ÅÅÅÅÇ (r + 3/2)Fr + 1/2(ξ) É 2Ñ ij (r + 5/2)Fr + 3/2(ξ) yz ÑÑÑÑ zz ÑÑ − jjjj zz ÑÑ k (r + 3/2)Fr + 1/2(ξ) { ÑÑÑÖ Fn(ξ) =

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time for phonon−phonon scattering (C is a constant and T is the temperature), which can be determined by κ pure = kB2 ΘD/(2π 2νsℏCT )

where ΘD is the Debye temperature. ℏ is the Planck constant. A is the coefficient can be given by

(3)

A = Ω 0Γ/(4πνs 3)

(8)

where Ω0 is the unit cell volume. Γ is the scattering parameter, which can be calculated as Γ = Γm + Γs, where Γm and Γs are scattering parameters related to mass fluctuation and strain field fluctuation, respectively. Γm and Γs can be expressed as follows53

(4)

(5)

Here, r is the scattering parameter, kB is the Boltzmann constant, e is the electron charge, Fn(ξ) is the Fermi integral, and ξ is the reduced Fermi energy. The calculated L values are shown in Figure 9b, which increases with increasing x, demonstrating an intensive degenerate degree. The obtained κe from κe = LσT in Figure 9c shows that the κe at 0 ≤ x ≤ 0.5 is low, which is no more than 0.2 W m−1 K−1. However, the κe increases sharply at x = 0.7 and x = 1, owing to the extremely high σ. We can thus estimate the κl by subtracting κe from κ (κl = κ − κe = κ − LσT), as shown in Figure 9d. The κl at x = 0 ranges from 0.42 to 0.45 W m−1 K−1 and has a decline at 0.02 ≤ x ≤ 0.1, which reflects intensive phonon point defect scattering owing to the mismatch of the size and mass by the substitution of Te for S.50 The lowest κl reaches 0.24−0.26 W m−1 K−1 at x = 0.08 in the whole measuring temperature region. As increasing x, the κl becomes higher at x = 0.3 and x = 0.5, which ascribes to the weakened phonon scattering deriving from the denser structure (Figure 3). It is worth noting the negative κl at x = 0.7 and x = 1, revealing the larger κe than κ. This indicates that the equation κe = LσT is not suitable for these two samples because of both contributions of charge carriers and mobile Cu ions to the electrical transport. The amended equation can be expressed as κe = L(σ − σi)T + κi, where σi and κi are the conductivity and the thermal conductivity of ions,51 respectively. However, the determination of the exact κe values at x = 0.7 and x = 1 is still a challenge as the σi is difficult to measure. It should be noted here that the measured direction of thermal transport is perpendicular to the pressure direction, which is different from that of electrical owing to the limit of the bulk size, and it may bring some deviation in the final ZT calculation for x = 1 bulk. Thus, we fabricated a larger sample of Φ 12 × 12 mm and cut it along perpendicular and parallel to the pressure direction. The measured TE properties along the same direction are shown in the Supporting Information (Figure S2). It was found that the σ perpendicular to the pressure direction is a little higher than that of the parallel direction. However, because of the lower κ of parallel direction compared with perpendicular direction, the ZT of the two directions is almost the same. In order to better understand the point defect scattering mechanism including the mass and the strain difference contribution between S and Te, the Callaway model is used to calculate the κl of Cu2S1−xTex compounds, which can be expressed as52 κl = kB/[4πvs(ACT )1/2 ]

(7)

i M −M 2 yz 1 ji M̅ zy jj zz x(1 − x)jjj 1 zz 3 jk M̿ z{ k M̅ { 2

Γm =

2

2 1 i M̅ y i r − r2 yz zz Γs = jjjj zzzz x(1 − x)εjjj 1 3 k M̿ { k r̅ {

(9)

2

(10)

where the M̅ , the M̿ , and the r ̅ can be given by M̅ = M1x + M 2(1 − x)

(11)

1 2 M̅ + M3 3 3

(12)

M̿ =

r ̅ = r1x + r2(1 − x)

(13)

where M1, M2, and M3 are the atomic mass of Te, S, and Cu, respectively. r1 and r2 are the atomic radius of Te and S, respectively, x is the Te content in one molecular. The parameter ε is usually estimated by fitting the experimental results, in which the method has been used in other work.54 Figure 9e shows the contributions of Γm and Γs with a function of x. It is clear that the Γm and the Γs increase as increasing x, peaking at 0.3 ≤ x ≤ 0.5, and then show a downtrend with increasing x. The larger Γm than Γs implies that the effect of mass fluctuation is more remarkable than the strain fluctuation to reduce the κl. The black line in Figure 9f shows the calculated κl based on the above equations. The experimental κl (red dots) at room temperature in our work well falls on the calculated line, expect the deviation of x = 0.3 and x = 0.5 samples, which can be ascribed to the increased density and the low transverse phonon velocity of Cu2S caused by the extreme softening of shear modes. By combining the measured PF and κ, the temperature dependence of TE figure-of-merit ZT for Cu2S1−xTex bulk samples is shown in Figure 10a. It is found that most 0 < x < 1 samples have higher ZT than the x = 0 and x = 1 samples, owing to the enhanced PF and/or reduced κ. A maximum ZT value 1.18 was achieved at 723 K for Cu2S0.94Te0.06, which is about 3.7 times that (0.32) of the pristine Cu2S. Figure 10b compares the ZT of our Cu2S (■) and Cu2S0.94Te0.06 (●) as well as reported high-performance Cu2S-based compounds from literatures.20,22−25,54,55 It is clear that Cu2S0.94Te0.06 obtained higher ZT at 723 K and a comparable ZT at 323− 573 K. Table 1 also lists the ZT values of other sulfide TE materials reported in the references for comparison.48,51,56−65 It is clear to see that the ZT of 1.18 obtained in Cu2S0.94Te0.06 is higher than most of the other advanced sulfides.

(6)

Here, kB is the Boltzmann constant. vs is the mean sound velocity and vs = 2393 m s−1 for Cu2S.20 CT is the relaxation 32208


Research Article


about 3.7 times that (0.32) of pristine Cu2S and 14.8 times that (0.08) of pristine Cu2Te.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.8b11300. XRD patterns of powder samples within 2θ range of 20°−50° for Cu2S1−xTex and temperature dependence of electrical conductivity σ, Seebeck coefficient α, and power factor α2σ, thermal conductivity κ, and ZT, which were measured along the same directions (PDF)

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Phone: +86-010-62334195. ORCID

Bo-Ping Zhang: 0000-0003-1712-6868 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS This work was supported by National Key R&D Program of China (grant no. 2018YFB0703603) and the National Natural Science Foundation of China (grant no. 11474176).

Figure 10. Temperature dependence of ZT for the Cu2S1−xTex bulk samples (a), and the Cu2S1−xTex (x = 0, 0.06) bulks as well as reported high-performance Cu2S-based compounds from literatures (b).

■

Table 1. ZT Value for the Cu2S0.94Te0.06 and Other Sulfide TE Materials material

ZT value

p/n

references

Cu2S0.94Te0.06 Na0.05Cu9S5 Sn0.995Na0.005S Cu12Sb4S12.7 CuCrS2.01 PbS + 3.0 at. % CdS Cu0.92Zn0.08FeS2 Cu0.005Bi2S3 AgBi3S5 + 0.33% Cl Bi/Bi2S3 Ce0.03Bi2S3 Ni-doped CoSbS In1.95Mg0.05S3

1.18 at 723 K 1.1 at 773 K 0.65 at 823 K 0.65 at 723 K 0.15 at 673 K 1.3 at 923 K 0.26 at 630 K 0.34 at 573 K 1.0 at 800 K 0.36 at 623 K 0.33 at 573 K 0.5 at 873 K 0.53 at 700 K

p p p p p p n n n n n n n

this work 51 56 57 58 59 48 60 61 62 63 64 65

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4. CONCLUSIONS In summary, we synthesized a series of Cu2S1−xTex (x = 0, 0.02, 0.04, 0.06, 0.08, 0.1, 0.3, 0.5, 0.7, and 1) alloys by a facile, rapid method of MA & SPS technique. The phase composition and crystal structure vary with x, which are the pure M-Cu2S at x = 0, M-Cu2S1−xTex solid solution at 0.02 ≤ x ≤ 0.06, HCu2S1−xTex solid solution at 0.08 ≤ x ≤ 0.1, coexisted H-Cu2S and H-Cu2Te at 0.3 ≤ x ≤ 0.7, and pure H-Cu2Te at x = 1. The band gap narrows with increasing Te in Cu2S, resulting in an increased n and enhanced σ. The PF rose up to 953.9 μW m−1 K−2 at 773 K for the Cu2S0.3Te0.7 sample, and the κ obtained the minimum 0.24−0.26 W m−1 K−1 at 323−773 K for Cu2S0.92Te0.08 bulk. Owing to a high PF (429.1 μW m−1 K−2) and a low κ (0.26 W m−1 K−1), a peak ZT of 1.18 was obtained at 723 K for the Cu2S0.94Te0.06 compound, which is 32209


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High Thermoelectric Figure of Merit Achieved in Cu2S1âxTex Alloys

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