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Low-temperature phase transition and transport properties of β-Cu2−xSe fabricated using hydrothermal synthesis and evacuating-and-encapsulating sintering
Journal of the European Ceramic Society 36 (2016) 2755–2760
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Journal of the European Ceramic Society journal homepage: www.elsevier.com/locate/jeurceramsoc
Low-temperature phase transition and transport properties of -Cu2−x Se fabricated using hydrothermal synthesis and evacuating-and-encapsulating sintering Ankam Bhaskar a , Chih-Hui Hu a , Ching-Lin Chang b , Chia-Jyi Liu a,∗ a b
Department of Physics, National Changhua University of Education, Changhua 500, Taiwan Department of Physics, Tamkang University, Tamsui, 251, Taiwan
a r t i c l e
i n f o
Article history: Received 4 February 2016 Received in revised form 1 April 2016 Accepted 19 April 2016 Available online 26 April 2016 Keywords: Phonon-liquid electron-crystal Cu2 Se Hydrothermal synthesis Transport properties
a b s t r a c t A series of phonon-liquid electron-crystal materials of -Cu2−x Se (0 ≤ x ≤ 0.04) is fabricated using hydrothermal synthesis followed by evacuating-and-encapsulating sintering. These materials are characterized using powder X-ray diffraction, field-emission scanning electron micrograph, electronic and thermal transport. Low-temperature X-ray diffraction patterns along with the electronic and thermal transport data reveal that -Cu2 Se exhibit a phase transition at around 160 K. We find that all the electrical conductivity, carrier concentration and thermopower of -Cu2−x Se increase with copper deficiency content x. A knee behavior is observed in the temperature dependence of both electrical resistivity and thermopower on cooling the samples from room temperature to around 160 K, which could be associated with the structural transition at the corresponding temperature. The thermal conductivity of -Cu2−x Se at room temperature is found to be larger than that of ␣-Cu2−x Se. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction P-type copper selenides Cu2 Se belongs to a member of transition metal chalcogenides, which have recently received great attention due to their intriguing properties and potential applications such as solar energy [1], gas sensors [2], optical filters [3], and supersonics [4]. Besides, it has been recently reported that a simple face-centered-cubic phase of Cu2 Se reaches a zT of 1.5 at 1000 K for thermoelectric energy conversion [5–7]. Due to the high diffusivity of copper ion in -Cu2 Se, this class of materials has been described as phonon-liquid electron-crystal system. Crystal packing of Cu2 Se differs at various temperatures to form cubic, tetragonal, orthorhombic, hexagonal and monoclinic polymorphs [8,9]. The designation for different polymorphs varies in literature. In this study, we adopt the monoclinic Cu2 Se as ␣ phase and the cubic Cu2 Se as  phase. The stoichiometric ␣-Cu2 Se is stable up to 414 K and undergoes a structural transition to high-temperature  phase [10,11]. For both of ␣ and  phases, the fcc lattice array is formed by selenium ions and the copper ions are distributed
∗ Corresponding author. E-mail address: [email protected] (C.-J. Liu). http://dx.doi.org/10.1016/j.jeurceramsoc.2016.04.031 0955-2219/© 2016 Elsevier Ltd. All rights reserved.
over tetrahedral 8c and 32f split positions within the octahedral voids (1/3 < x <1/2) [12]. However, it is controversial for the copper ion distribution at different interstitial positions [12–15]. Heyding et al. [14] reported that the copper ions were located at tetrahedral 8c (1/4,1/4, 1/4) positions and the remaining Cu ions were at trigonal 32f (x, x, x) with x = 0.33 positions. Skomorokhov et al. [12] reported that the copper ions was at trigonal 32f1 (x = 0.297) and 32f2 (x = 0.471) positions. Yamamoto et al. [13] reported that the copper ions were located at tetrahedral and trigonal positions and a small fraction of copper ions at octahedral 4b (1/2, 1/2, 1/2) positions. Since the copper ions jump between available lattice sites within the rigid fcc selenium cage at high temperatures, the Cu2 Se with high mobility of copper ions is considered as superionic phase with larger ionic conductivity than the ␣-phase. Various fabrication techniques such as solid state reaction [10–12], sonochemical synthesis [13], sol-gel, [3] and microwaveassisted heating [14] have been adopted to prepare Cu2 Se. The solid state reaction is simple but requires high temperature treatment with long soaking time to obtain pure Cu2 Se phase [10–12]. The major drawback of microwave-assisted heating is the high investment cost and short penetration depth of microwave irradiation into the specimen. Hydrothermal synthesis can avoid these problems and can be considered as a promising way for synthesizing
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nanocrystalline materials. Besides, there are several advantages such as simplicity, products with high crystallinity, relatively low reaction temperature, and control of morphology. To the best of our knowledge, there are two reports regarding the Cu2 Se powders synthesized using hydrothermal co-reduction method, which only report X-ray diffraction patterns and surface morphology of samples without transport properties measured [15,16]. Fabrication copper selenide using solid state reaction often leads to formation of ␣ phase; nevertheless hydrothermal synthesis leads to  phase. Furthermore, it has been suggested that the structural transition from ␣ phase to  phase is accompanied by a release of copper, which would then be oxidized to copper oxides [13]. There is no report so far on the transport properties of -Cu2−x Se fabricated using hydrothermal method. Motivated by this situation, we investigate low temperature structural and transport properties of hydrothermally synthesized -Cu2−x Se which is then consolidated by evacuated-and-encapsulated sintering. We find that all the electrical conductivity, carrier concentration and thermopower of -Cu2−x Se increase with copper deficiency content x. Besides, a knee behavior is observed in the temperature dependence of both electrical resistivity and thermopower on cooling the samples from room temperature to around 160 K, which could be associated with the structural transition at the corresponding temperature. The thermal conductivity of -Cu2−x Se at room temperature is found to be larger than that of ␣-Cu2−x Se.
2. Experimental procedure The reactants of CuCl2 ·2H2 O and Se powders were weighted in an appropriate ratio to have a nominal composition of Cu2−x Se (0 ≤ x ≤ 0.04). Sodium borohydride (NaBH4 ) was used as reducing agent. A sufficient amount of NaBH4 was slowly added into the Teflon bottle containing aqueous CuCl2 solution. The mixed solution was stirred until no gas bubbling, the autoclave was then sealed immediately. The autoclave is heated to 200 ◦ C with a ramp rate of 1 ◦ C/min and kept for 12 h. After heating, it was cooled down to room temperature naturally. The resulting product was filtered using suction filtration, washed several times with distilled water and ethanol, and then dried at 80 ◦ C. The dark-green powders were obtained and then cold pressed into a parallelepiped. The resulting parallelepipeds were then loaded into a Pyrex ampoule, which were evacuated using a diffusion pump to reach 10−5 Torr to 10−6 Torr and then sealed. The parallelepipeds in the encapsulated ampoule were then sintered at three different temperatures of 250 ◦ C, 350 ◦ C and 450 ◦ C for 12 h in three batches. The phase purity of the resulting powders was examined by a Shimadzu XRD-6000 powder X-ray diffractometer equipped with Fe K˛ radiation. Low-temperature powder X-ray diffraction patterns were obtained with Cu K˛ radiation. The morphology of the samples was examined using a Zeiss AURIGA field emission scanning electron microscope (FE-SEM). Both of the electrical resistivity and thermopower measurements were performed on cooling from room temperature to 80 K. Electrical resistivity and thermopower measurements were carried out using standard four-probe techniques and steady-state techniques, respectively. A type E differential thermocouple was used to measure the temperature difference between hot and cold ends of the sample [18], which was measured using a Keithley 2182 nanovoltmeter. The thermopower of the sample was obtained by subtracting the thermopower of the Cu Seebeck probe. Carrier concentration and mobility were determined by the Hall measurements using the van der Pauw method under an applied field of 0.55 T (ECOPIA: HMS-3000). Thermal conductivity measurements were carried out using transient plane source techniques with very small temperature perturbations of the sample using a Hot Disk thermal constants analyzer [19,20].
Fig. 1. XRD patterns of as-synthesized -Cu2−x Se with x = 0.00, 0.02 and 0.04. The ‘x’ and ‘o’ marks indicate the impurities of Cu2 O and CuO, respectively.
The transient plane source technique makes use of a thin sensor element in the shape of a double spiral. The Hot Disk sensor acts both as a heat source for generating a temperature gradient in the sample and a resistance thermometer for recording the time dependent temperature increase. A Kapton-insulated sensor was sandwiched between two pieces of samples. During a preset time, 200 resistance recordings were taken, and from these a relation between the temperature and time was established. The uncertainty for the electrical resistivity, thermopower, and thermal conductivity is about ±3%, ±4% and ±4%, respectively. The chemical composition is determined using a JEOL JXA-8200 Electron Probe X-Ray Microanalyzer (EPMA). The relative bulk density of all the samples was measured using the Archimedes’ method with an uncertainty of about ±2%. 3. Results and discussion Fig. 1 shows the X-ray diffraction patterns (XRD) of hydrothermally synthesized -Cu2−x Se with x = 0.00, 0.02 and 0.04. In addition to impurities of CuO (JCPDS: 80-1917) and Cu2 O (JCPDS: 78-2076), the XRD patterns show that all the samples can be indexed based on a -Cu2 Se cubic structure (JCPDS: 6-680) with space group # 225 of Fm3m. Fig. 2 shows the XRD patterns of -Cu2 Se sintered at 250 ◦ C, 350 ◦ C and 450 ◦ C, respectively, and Cu2−x Se with x = 0.02 and 0.04 sintered at 350 ◦ C. The XRD patterns clearly indicate that the impurities of CuO and Cu2 O in the assynthesized powders are drastically reduced after sintering. When the sintering temperature is increased from 350 ◦ C to 450 ◦ C, the XRD pattern indicates that a small portion of the  phase converts to the ␣ phase with monoclinic structure (JCPDS: 27-1131). This is in contrast to earlier reports in that the crystalline phase is stable as ␣-Cu2 Se at low temperatures and -Cu2 Se at high temperatures with its structural phase transition occurring at the temperature between 415 K and 423 K [10–12]. Furthermore, it has been reported that copper is released while the phase transition (␣ → ) occurs, and then Cu2 O is formed [11]. The refined lattice constant is a = 5.76(8) Å, 5.78(1) Å, 5.78(1) Å for x = 0.00 sintered at 250 ◦ C, 350 ◦ C and 450 ◦ C, respectively; the lattice constant is 5.78(6) Å, 5.78(7) Å for Cu2−x Se with x = 0.02 and 0.04 sintered at 350 ◦ C, respectively. These results are in relatively good agreement with previous reports [2,7,11,12,21]. The EPMA results indicate that the atomic ratio of Cu and Se is 2.03 and 1.98 for the x = 0.00 and 0.02 samples, respectively. The density of the samples falls in the range of 79 and 80%.
A. Bhaskar et al. / Journal of the European Ceramic Society 36 (2016) 2755–2760
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Fig. 2. XRD patterns of -Cu2 Se sintered at 250 ◦ C, 350◦ C and 450◦ C, and -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C. The ‘x’ and ‘o’ marks indicate the impurities of Cu2 O and CuO, respectively. The red and purple Miller indexes indicate the reflection planes of ␣- and -Cu2 Se, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 3 shows the FE-SEM images of fractured surface of Cu2−x Se with x = 0.00, 0.02 and 0.04. The layer morphology of -Cu2−x Se can be clearly observed. Various sizes of nanoplates are observed for the x = 0.00 sample. The surface morphology of the x = 0.02 and 0.04 samples shows leaf-like nanoplates. Similar morphology was found in -Cu2 Se synthesized using solvothermal method [22,23]. Fig. 4 shows the temperature dependence of electrical resistivity () for -Cu2 Se sintered at 250 ◦ C, 350 ◦ C and 450 ◦ C, and -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350 ◦ C. The magnitude of for all the samples is relatively low and in the range between 0.1 m cm and 0.6 m cm. The electrical resistivity decreases with increasing sintering temperature. The electrical resistivity drastically reduces for the sample sintered at 450 ◦ C. The electrical resistivity decreases with increasing deficiency of copper content due to an increase of hole carrier concentration. According to earlier reports [10,12], the deficiency of copper ions generates hole carriers into the -Cu2−x Se system, which in turn will decrease the electrical resistivity. Hall measurements confirm that carrier concentration increases with increasing deficiency of copper content. The carrier concentration and mobility at room temperature are 5.6 × 1019 cm−3 and 886 cm2 /Vs for x = 0.00, 6.0 × 1019 cm−3 and 622 cm2 /Vs for x = 0.02, 1.0 × 1020 cm−3 and 583 cm2 /Vs for x = 0.04, respectively. The mobility of hydrothermally synthesized -Cu2−x Se is significantly larger than that (14–20 cm2 /Vs) of ␣Cu2−x Se in Ref. [4]. The electrical resistivity of all the samples increases with increasing temperature. As shown in Fig. 4, the electrical resistivity has a sudden change at 162 K for the Cu2 Se samples sintered at 250 ◦ C, 350 ◦ C and 450 ◦ C, respectively. This could be associated with the structural transition at the corresponding temperatures. Fig. 5 shows the XRD patterns of the -Cu2 Se sample sintered at 250 ◦ C measured at 100, 150, 200 and 298 K, respectively, which are a zoomed view of the twotheta region 23◦ ≤ 2 ≤ 30◦ of the same patterns shown in the insert with 10◦ ≤ 2 ≤ 90◦ . As the temperature cooled to 200 K, the peaks occurring at 2 = 13.25◦ and 25.61◦ belong to the (030) and (060) reflections of the ␣-phase, respectively, indicative of a small amount of -phase transforming to ␣-phase. Upon further cooling to 150 K and below, two additional peaks appear at 2 = 24.28◦ and 25.55◦ , which has been ascribed to the formation of a superstructure of the -phase [12]. Fig. 6 shows the temperature dependence
Fig. 3. FE-SEM images of fractured surface of -Cu2−x Se (a) with x = 0.00 sintered at 250 ◦ C; (b) 0.02; and (c) 0.04 sintered at 350 ◦ C.
Fig. 4. The temperature dependence of electrical resistivity for (a) -Cu2−x Se sintered at various temperatures; (b) -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C.
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Fig. 5. The XRD patterns showing the additional peaks occurring at 2 = 23−30◦ , which could be associated with the superstructure of the -Cu2−x Se phase. The inset shows the XRD patterns of -Cu2 Se sintered at 250 ◦ C with 10◦ ≤ 2 ≤ 90◦ at various heating temperatures. The incident X-ray for the low temperature measurements is Cu K␣ radiation.
Fig. 6. The temperature dependence of electrical resistivity for -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350 ◦ C.
of electrical resistivity for -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C. The electrical resistivity also shows an abrupt change at ca. 162 K for x = 0.00 and at ca. 157 K for x = 0.02 and 0.04, respectively. Chi et al. [17] reported an anomalous behavior of resistivity for monoclinic Cu2 Se between 100 and 150 K, which was ascribed to a reversible superstructure transition of ␣-phase.
It should be noted that a resistivity hump is observed between 100 and 150 K in Ref. [25], which is obviously different from our resistivity data distinctly showing a “knee” around 160 K. Furthermore, the anomalous hump behavior in electrical resistivity in monoclinic Cu2 Se of Ref. [17] resembles that for -Cu2−x Se of Ref. [12] during cooling cycle measurements. Thermopower measurements are a very sensitive probe to the type and characteristic energy of carriers and are a complementary tool to the resistivity measurements for transport studies. Since thermopower is a measure of the heat per carrier over temperature, we can thus view it as a measure of the entropy per carrier. The positive sign of thermopower for all the samples indicates that holes are the dominant carriers. Hall measurements also confirm that the majority carriers are holes. Fig. 7 shows the temperature dependence of thermopower (S) for -Cu2 Se sintered at 250 ◦ C, 350 ◦ C and 450 ◦ C, and -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350 ◦ C. It can be readily seen that the size of thermopower increases with increasing temperature and deficiency content of copper. The thermopower also exhibits a “knee” temperature dependence with the bending angle depending on the sintering temperature and deficiency of copper. By comparison, Chi et al. [17] stated that there was no such a “knee” anomaly found in thermopower measurements. Based on the high electrical conductivity, carrier concentration, and small thermopower of -Cu2 Se can be considered as a highly degenerate
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Fig. 8. The temperature dependence of power factor for (a) -Cu2 Se sintered at various temperatures; (b) -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C. Fig. 7. The temperature dependence of thermopower for (a) -Cu2 Se sintered at various temperatures; (b) -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C.
semiconductor [8], thermopower can therefore be expressed as [24–26], S(T ) =
82 kB2 T 3eh2
m∗ T
2/3 3n
(1)
where n is the carrier concentration and m* the effective mass of the carrier. Nevertheless, based on the Hall concentration and mobility data for -Cu2−x Se, it seems that Eq. (1) is inadequate to describe the variation of thermopower with the deficiency content of x since the carrier concentration increases and the mobility decreases with increasing x. The power factor (P = S2 /) is calculated from the measured thermopower and electrical resistivity. Fig. 8 shows the temperature dependence of thermoelectric power factor (S2 /) for -Cu2 Se sintered at 250◦ C, 350◦ C and 450◦ C, and -Cu2−x Se with x = 0.00, 0.02 and 0.04 samples sintered at 350◦ C. Among the samples sintered at 350◦ C, -Cu1.96 Se has the highest power factor. Fig. 9 shows the thermal conductivity of -Cu2 Se sintered at 250◦ C, 350◦ C and 450◦ C, and -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C. The thermal conductivity of -Cu2−x Se at room temperature is found to be larger than that of ␣-Cu2−x Se [5]. Accompanying the ␣ to  phase transition for Cu2 Se fabricated by solid state reaction, a discontinuous increase of thermal conductivity is observed. The high diffusivity of copper ion has been described as a diffusion of “liquid-like” charged fluid of coper ions in a crystalline selenium sublattice [5]. The total thermal conduc-
tivity shows a decreasing trend with increasing temperature and a plateau behavior between 160 and 190 K. The plateau behavior should be associated with the structural phase transition. Xing et al. reported that the total thermal conductivity is in the range between 2.5 W m−1 K−1 and 6 W m−1 K−1 at room temperature for Cu2−x Se with x = 0.00 to 0.25 [10]. The calculated zT value is found to be very small but increases with increasing temperature. Cu1.96 Se shows the highest zT = 0.033 at 260 K among the samples. 4. Conclusions We have fabricated -Cu2−x Se with x = 0.00, 0.02, and 0.04 using hydrothermal synthesis followed by evacuating-andencapsulating sintering. All the electrical conductivity, carrier concentration and thermopower of -Cu2−x Se increase with copper deficiency content x. The thermal conductivity of -Cu2−x Se at room temperature is found to be larger than that of ␣-Cu2−x Se. Besides, a knee behavior is observed in the temperature dependence of both electrical resistivity and thermopower on cooling the samples from room temperature to around 160 K, which could be associated with the structural transition at the corresponding temperature. Acknowledgments This work was supported by Ministry of Science and Technology of Taiwan under the Grant No. 101-2112-M-018-003-MY3 and 104-2112-M-018-002-MY3. Ankam Bhaskar would like to
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Fig. 9. The temperature dependence of total thermal conductivity for (a) -Cu2 Se sintered at various temperatures; (b) -Cu2-x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C.
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Contents lists available at www.sciencedirect.com
Journal of the European Ceramic Society journal homepage: www.elsevier.com/locate/jeurceramsoc
Low-temperature phase transition and transport properties of -Cu2−x Se fabricated using hydrothermal synthesis and evacuating-and-encapsulating sintering Ankam Bhaskar a , Chih-Hui Hu a , Ching-Lin Chang b , Chia-Jyi Liu a,∗ a b
Department of Physics, National Changhua University of Education, Changhua 500, Taiwan Department of Physics, Tamkang University, Tamsui, 251, Taiwan
a r t i c l e
i n f o
Article history: Received 4 February 2016 Received in revised form 1 April 2016 Accepted 19 April 2016 Available online 26 April 2016 Keywords: Phonon-liquid electron-crystal Cu2 Se Hydrothermal synthesis Transport properties
a b s t r a c t A series of phonon-liquid electron-crystal materials of -Cu2−x Se (0 ≤ x ≤ 0.04) is fabricated using hydrothermal synthesis followed by evacuating-and-encapsulating sintering. These materials are characterized using powder X-ray diffraction, field-emission scanning electron micrograph, electronic and thermal transport. Low-temperature X-ray diffraction patterns along with the electronic and thermal transport data reveal that -Cu2 Se exhibit a phase transition at around 160 K. We find that all the electrical conductivity, carrier concentration and thermopower of -Cu2−x Se increase with copper deficiency content x. A knee behavior is observed in the temperature dependence of both electrical resistivity and thermopower on cooling the samples from room temperature to around 160 K, which could be associated with the structural transition at the corresponding temperature. The thermal conductivity of -Cu2−x Se at room temperature is found to be larger than that of ␣-Cu2−x Se. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction P-type copper selenides Cu2 Se belongs to a member of transition metal chalcogenides, which have recently received great attention due to their intriguing properties and potential applications such as solar energy [1], gas sensors [2], optical filters [3], and supersonics [4]. Besides, it has been recently reported that a simple face-centered-cubic phase of Cu2 Se reaches a zT of 1.5 at 1000 K for thermoelectric energy conversion [5–7]. Due to the high diffusivity of copper ion in -Cu2 Se, this class of materials has been described as phonon-liquid electron-crystal system. Crystal packing of Cu2 Se differs at various temperatures to form cubic, tetragonal, orthorhombic, hexagonal and monoclinic polymorphs [8,9]. The designation for different polymorphs varies in literature. In this study, we adopt the monoclinic Cu2 Se as ␣ phase and the cubic Cu2 Se as  phase. The stoichiometric ␣-Cu2 Se is stable up to 414 K and undergoes a structural transition to high-temperature  phase [10,11]. For both of ␣ and  phases, the fcc lattice array is formed by selenium ions and the copper ions are distributed
∗ Corresponding author. E-mail address: [email protected] (C.-J. Liu). http://dx.doi.org/10.1016/j.jeurceramsoc.2016.04.031 0955-2219/© 2016 Elsevier Ltd. All rights reserved.
over tetrahedral 8c and 32f split positions within the octahedral voids (1/3 < x <1/2) [12]. However, it is controversial for the copper ion distribution at different interstitial positions [12–15]. Heyding et al. [14] reported that the copper ions were located at tetrahedral 8c (1/4,1/4, 1/4) positions and the remaining Cu ions were at trigonal 32f (x, x, x) with x = 0.33 positions. Skomorokhov et al. [12] reported that the copper ions was at trigonal 32f1 (x = 0.297) and 32f2 (x = 0.471) positions. Yamamoto et al. [13] reported that the copper ions were located at tetrahedral and trigonal positions and a small fraction of copper ions at octahedral 4b (1/2, 1/2, 1/2) positions. Since the copper ions jump between available lattice sites within the rigid fcc selenium cage at high temperatures, the Cu2 Se with high mobility of copper ions is considered as superionic phase with larger ionic conductivity than the ␣-phase. Various fabrication techniques such as solid state reaction [10–12], sonochemical synthesis [13], sol-gel, [3] and microwaveassisted heating [14] have been adopted to prepare Cu2 Se. The solid state reaction is simple but requires high temperature treatment with long soaking time to obtain pure Cu2 Se phase [10–12]. The major drawback of microwave-assisted heating is the high investment cost and short penetration depth of microwave irradiation into the specimen. Hydrothermal synthesis can avoid these problems and can be considered as a promising way for synthesizing
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nanocrystalline materials. Besides, there are several advantages such as simplicity, products with high crystallinity, relatively low reaction temperature, and control of morphology. To the best of our knowledge, there are two reports regarding the Cu2 Se powders synthesized using hydrothermal co-reduction method, which only report X-ray diffraction patterns and surface morphology of samples without transport properties measured [15,16]. Fabrication copper selenide using solid state reaction often leads to formation of ␣ phase; nevertheless hydrothermal synthesis leads to  phase. Furthermore, it has been suggested that the structural transition from ␣ phase to  phase is accompanied by a release of copper, which would then be oxidized to copper oxides [13]. There is no report so far on the transport properties of -Cu2−x Se fabricated using hydrothermal method. Motivated by this situation, we investigate low temperature structural and transport properties of hydrothermally synthesized -Cu2−x Se which is then consolidated by evacuated-and-encapsulated sintering. We find that all the electrical conductivity, carrier concentration and thermopower of -Cu2−x Se increase with copper deficiency content x. Besides, a knee behavior is observed in the temperature dependence of both electrical resistivity and thermopower on cooling the samples from room temperature to around 160 K, which could be associated with the structural transition at the corresponding temperature. The thermal conductivity of -Cu2−x Se at room temperature is found to be larger than that of ␣-Cu2−x Se.
2. Experimental procedure The reactants of CuCl2 ·2H2 O and Se powders were weighted in an appropriate ratio to have a nominal composition of Cu2−x Se (0 ≤ x ≤ 0.04). Sodium borohydride (NaBH4 ) was used as reducing agent. A sufficient amount of NaBH4 was slowly added into the Teflon bottle containing aqueous CuCl2 solution. The mixed solution was stirred until no gas bubbling, the autoclave was then sealed immediately. The autoclave is heated to 200 ◦ C with a ramp rate of 1 ◦ C/min and kept for 12 h. After heating, it was cooled down to room temperature naturally. The resulting product was filtered using suction filtration, washed several times with distilled water and ethanol, and then dried at 80 ◦ C. The dark-green powders were obtained and then cold pressed into a parallelepiped. The resulting parallelepipeds were then loaded into a Pyrex ampoule, which were evacuated using a diffusion pump to reach 10−5 Torr to 10−6 Torr and then sealed. The parallelepipeds in the encapsulated ampoule were then sintered at three different temperatures of 250 ◦ C, 350 ◦ C and 450 ◦ C for 12 h in three batches. The phase purity of the resulting powders was examined by a Shimadzu XRD-6000 powder X-ray diffractometer equipped with Fe K˛ radiation. Low-temperature powder X-ray diffraction patterns were obtained with Cu K˛ radiation. The morphology of the samples was examined using a Zeiss AURIGA field emission scanning electron microscope (FE-SEM). Both of the electrical resistivity and thermopower measurements were performed on cooling from room temperature to 80 K. Electrical resistivity and thermopower measurements were carried out using standard four-probe techniques and steady-state techniques, respectively. A type E differential thermocouple was used to measure the temperature difference between hot and cold ends of the sample [18], which was measured using a Keithley 2182 nanovoltmeter. The thermopower of the sample was obtained by subtracting the thermopower of the Cu Seebeck probe. Carrier concentration and mobility were determined by the Hall measurements using the van der Pauw method under an applied field of 0.55 T (ECOPIA: HMS-3000). Thermal conductivity measurements were carried out using transient plane source techniques with very small temperature perturbations of the sample using a Hot Disk thermal constants analyzer [19,20].
Fig. 1. XRD patterns of as-synthesized -Cu2−x Se with x = 0.00, 0.02 and 0.04. The ‘x’ and ‘o’ marks indicate the impurities of Cu2 O and CuO, respectively.
The transient plane source technique makes use of a thin sensor element in the shape of a double spiral. The Hot Disk sensor acts both as a heat source for generating a temperature gradient in the sample and a resistance thermometer for recording the time dependent temperature increase. A Kapton-insulated sensor was sandwiched between two pieces of samples. During a preset time, 200 resistance recordings were taken, and from these a relation between the temperature and time was established. The uncertainty for the electrical resistivity, thermopower, and thermal conductivity is about ±3%, ±4% and ±4%, respectively. The chemical composition is determined using a JEOL JXA-8200 Electron Probe X-Ray Microanalyzer (EPMA). The relative bulk density of all the samples was measured using the Archimedes’ method with an uncertainty of about ±2%. 3. Results and discussion Fig. 1 shows the X-ray diffraction patterns (XRD) of hydrothermally synthesized -Cu2−x Se with x = 0.00, 0.02 and 0.04. In addition to impurities of CuO (JCPDS: 80-1917) and Cu2 O (JCPDS: 78-2076), the XRD patterns show that all the samples can be indexed based on a -Cu2 Se cubic structure (JCPDS: 6-680) with space group # 225 of Fm3m. Fig. 2 shows the XRD patterns of -Cu2 Se sintered at 250 ◦ C, 350 ◦ C and 450 ◦ C, respectively, and Cu2−x Se with x = 0.02 and 0.04 sintered at 350 ◦ C. The XRD patterns clearly indicate that the impurities of CuO and Cu2 O in the assynthesized powders are drastically reduced after sintering. When the sintering temperature is increased from 350 ◦ C to 450 ◦ C, the XRD pattern indicates that a small portion of the  phase converts to the ␣ phase with monoclinic structure (JCPDS: 27-1131). This is in contrast to earlier reports in that the crystalline phase is stable as ␣-Cu2 Se at low temperatures and -Cu2 Se at high temperatures with its structural phase transition occurring at the temperature between 415 K and 423 K [10–12]. Furthermore, it has been reported that copper is released while the phase transition (␣ → ) occurs, and then Cu2 O is formed [11]. The refined lattice constant is a = 5.76(8) Å, 5.78(1) Å, 5.78(1) Å for x = 0.00 sintered at 250 ◦ C, 350 ◦ C and 450 ◦ C, respectively; the lattice constant is 5.78(6) Å, 5.78(7) Å for Cu2−x Se with x = 0.02 and 0.04 sintered at 350 ◦ C, respectively. These results are in relatively good agreement with previous reports [2,7,11,12,21]. The EPMA results indicate that the atomic ratio of Cu and Se is 2.03 and 1.98 for the x = 0.00 and 0.02 samples, respectively. The density of the samples falls in the range of 79 and 80%.
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Fig. 2. XRD patterns of -Cu2 Se sintered at 250 ◦ C, 350◦ C and 450◦ C, and -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C. The ‘x’ and ‘o’ marks indicate the impurities of Cu2 O and CuO, respectively. The red and purple Miller indexes indicate the reflection planes of ␣- and -Cu2 Se, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
Fig. 3 shows the FE-SEM images of fractured surface of Cu2−x Se with x = 0.00, 0.02 and 0.04. The layer morphology of -Cu2−x Se can be clearly observed. Various sizes of nanoplates are observed for the x = 0.00 sample. The surface morphology of the x = 0.02 and 0.04 samples shows leaf-like nanoplates. Similar morphology was found in -Cu2 Se synthesized using solvothermal method [22,23]. Fig. 4 shows the temperature dependence of electrical resistivity () for -Cu2 Se sintered at 250 ◦ C, 350 ◦ C and 450 ◦ C, and -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350 ◦ C. The magnitude of for all the samples is relatively low and in the range between 0.1 m cm and 0.6 m cm. The electrical resistivity decreases with increasing sintering temperature. The electrical resistivity drastically reduces for the sample sintered at 450 ◦ C. The electrical resistivity decreases with increasing deficiency of copper content due to an increase of hole carrier concentration. According to earlier reports [10,12], the deficiency of copper ions generates hole carriers into the -Cu2−x Se system, which in turn will decrease the electrical resistivity. Hall measurements confirm that carrier concentration increases with increasing deficiency of copper content. The carrier concentration and mobility at room temperature are 5.6 × 1019 cm−3 and 886 cm2 /Vs for x = 0.00, 6.0 × 1019 cm−3 and 622 cm2 /Vs for x = 0.02, 1.0 × 1020 cm−3 and 583 cm2 /Vs for x = 0.04, respectively. The mobility of hydrothermally synthesized -Cu2−x Se is significantly larger than that (14–20 cm2 /Vs) of ␣Cu2−x Se in Ref. [4]. The electrical resistivity of all the samples increases with increasing temperature. As shown in Fig. 4, the electrical resistivity has a sudden change at 162 K for the Cu2 Se samples sintered at 250 ◦ C, 350 ◦ C and 450 ◦ C, respectively. This could be associated with the structural transition at the corresponding temperatures. Fig. 5 shows the XRD patterns of the -Cu2 Se sample sintered at 250 ◦ C measured at 100, 150, 200 and 298 K, respectively, which are a zoomed view of the twotheta region 23◦ ≤ 2 ≤ 30◦ of the same patterns shown in the insert with 10◦ ≤ 2 ≤ 90◦ . As the temperature cooled to 200 K, the peaks occurring at 2 = 13.25◦ and 25.61◦ belong to the (030) and (060) reflections of the ␣-phase, respectively, indicative of a small amount of -phase transforming to ␣-phase. Upon further cooling to 150 K and below, two additional peaks appear at 2 = 24.28◦ and 25.55◦ , which has been ascribed to the formation of a superstructure of the -phase [12]. Fig. 6 shows the temperature dependence
Fig. 3. FE-SEM images of fractured surface of -Cu2−x Se (a) with x = 0.00 sintered at 250 ◦ C; (b) 0.02; and (c) 0.04 sintered at 350 ◦ C.
Fig. 4. The temperature dependence of electrical resistivity for (a) -Cu2−x Se sintered at various temperatures; (b) -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C.
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Fig. 5. The XRD patterns showing the additional peaks occurring at 2 = 23−30◦ , which could be associated with the superstructure of the -Cu2−x Se phase. The inset shows the XRD patterns of -Cu2 Se sintered at 250 ◦ C with 10◦ ≤ 2 ≤ 90◦ at various heating temperatures. The incident X-ray for the low temperature measurements is Cu K␣ radiation.
Fig. 6. The temperature dependence of electrical resistivity for -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350 ◦ C.
of electrical resistivity for -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C. The electrical resistivity also shows an abrupt change at ca. 162 K for x = 0.00 and at ca. 157 K for x = 0.02 and 0.04, respectively. Chi et al. [17] reported an anomalous behavior of resistivity for monoclinic Cu2 Se between 100 and 150 K, which was ascribed to a reversible superstructure transition of ␣-phase.
It should be noted that a resistivity hump is observed between 100 and 150 K in Ref. [25], which is obviously different from our resistivity data distinctly showing a “knee” around 160 K. Furthermore, the anomalous hump behavior in electrical resistivity in monoclinic Cu2 Se of Ref. [17] resembles that for -Cu2−x Se of Ref. [12] during cooling cycle measurements. Thermopower measurements are a very sensitive probe to the type and characteristic energy of carriers and are a complementary tool to the resistivity measurements for transport studies. Since thermopower is a measure of the heat per carrier over temperature, we can thus view it as a measure of the entropy per carrier. The positive sign of thermopower for all the samples indicates that holes are the dominant carriers. Hall measurements also confirm that the majority carriers are holes. Fig. 7 shows the temperature dependence of thermopower (S) for -Cu2 Se sintered at 250 ◦ C, 350 ◦ C and 450 ◦ C, and -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350 ◦ C. It can be readily seen that the size of thermopower increases with increasing temperature and deficiency content of copper. The thermopower also exhibits a “knee” temperature dependence with the bending angle depending on the sintering temperature and deficiency of copper. By comparison, Chi et al. [17] stated that there was no such a “knee” anomaly found in thermopower measurements. Based on the high electrical conductivity, carrier concentration, and small thermopower of -Cu2 Se can be considered as a highly degenerate
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Fig. 8. The temperature dependence of power factor for (a) -Cu2 Se sintered at various temperatures; (b) -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C. Fig. 7. The temperature dependence of thermopower for (a) -Cu2 Se sintered at various temperatures; (b) -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C.
semiconductor [8], thermopower can therefore be expressed as [24–26], S(T ) =
82 kB2 T 3eh2
m∗ T
2/3 3n
(1)
where n is the carrier concentration and m* the effective mass of the carrier. Nevertheless, based on the Hall concentration and mobility data for -Cu2−x Se, it seems that Eq. (1) is inadequate to describe the variation of thermopower with the deficiency content of x since the carrier concentration increases and the mobility decreases with increasing x. The power factor (P = S2 /) is calculated from the measured thermopower and electrical resistivity. Fig. 8 shows the temperature dependence of thermoelectric power factor (S2 /) for -Cu2 Se sintered at 250◦ C, 350◦ C and 450◦ C, and -Cu2−x Se with x = 0.00, 0.02 and 0.04 samples sintered at 350◦ C. Among the samples sintered at 350◦ C, -Cu1.96 Se has the highest power factor. Fig. 9 shows the thermal conductivity of -Cu2 Se sintered at 250◦ C, 350◦ C and 450◦ C, and -Cu2−x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C. The thermal conductivity of -Cu2−x Se at room temperature is found to be larger than that of ␣-Cu2−x Se [5]. Accompanying the ␣ to  phase transition for Cu2 Se fabricated by solid state reaction, a discontinuous increase of thermal conductivity is observed. The high diffusivity of copper ion has been described as a diffusion of “liquid-like” charged fluid of coper ions in a crystalline selenium sublattice [5]. The total thermal conduc-
tivity shows a decreasing trend with increasing temperature and a plateau behavior between 160 and 190 K. The plateau behavior should be associated with the structural phase transition. Xing et al. reported that the total thermal conductivity is in the range between 2.5 W m−1 K−1 and 6 W m−1 K−1 at room temperature for Cu2−x Se with x = 0.00 to 0.25 [10]. The calculated zT value is found to be very small but increases with increasing temperature. Cu1.96 Se shows the highest zT = 0.033 at 260 K among the samples. 4. Conclusions We have fabricated -Cu2−x Se with x = 0.00, 0.02, and 0.04 using hydrothermal synthesis followed by evacuating-andencapsulating sintering. All the electrical conductivity, carrier concentration and thermopower of -Cu2−x Se increase with copper deficiency content x. The thermal conductivity of -Cu2−x Se at room temperature is found to be larger than that of ␣-Cu2−x Se. Besides, a knee behavior is observed in the temperature dependence of both electrical resistivity and thermopower on cooling the samples from room temperature to around 160 K, which could be associated with the structural transition at the corresponding temperature. Acknowledgments This work was supported by Ministry of Science and Technology of Taiwan under the Grant No. 101-2112-M-018-003-MY3 and 104-2112-M-018-002-MY3. Ankam Bhaskar would like to
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Fig. 9. The temperature dependence of total thermal conductivity for (a) -Cu2 Se sintered at various temperatures; (b) -Cu2-x Se with x = 0.00, 0.02 and 0.04 sintered at 350◦ C.
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