Ultra-low thermal conductivity and high thermoelectric performance realized in a Cu₃SbSe₄ based system

Abstract

Cu₃SbSe₄-Based materials were fabricated through Sn-doping and AgSb_0.98Ge_0.02Se₂ incorporation and their thermoelectric properties were investigated in the temperature range from 300 K to 675 K. A remarkable enhancement in thermoelectric performance was obtained, which can be ascribed to the synergistic adjustment of electrical and thermal transport. The introduction of AgSb_0.98Ge_0.02Se₂ nanoparticles strengthened phonon scattering in the composites, resulting in a low thermal conductivity. Simultaneously, a high power factor was maintained. As a result, a maximum figure of merit of 1.23 was obtained at 675 K for the sample Cu₃Sb_0.96Sn_0.04Se₄–3 wt% AgSb_0.98Ge_0.02Se₂, which is one of the highest values ever reported in the Cu₃SbSe₄-based system.

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1. Introduction

Energy shortages and environmental degradation have become major issues that have increasingly haunted humanity with the development of industrial society.¹ In this context, thermoelectric (TE) materials can be used as solid state refrigerators or heat pumps that can directly convert waste heat from industrial production into electricity and provide possible solutions to address energy shortages.^2,3 It is known that the TE energy conversion efficiency is defined by the dimensionless figure-of-merit ZT = S²σT/κ, where S, σ, T and κ are the Seebeck coefficient, electrical conductivity, absolute temperature and thermal conductivity, respectively.⁴ As one of the novel chalcogenide thermoelectric materials, the ternary cooper-based p-type semiconductor Cu₃SbSe₄ has attracted more attention due to its relatively low thermal conductivity and comparatively narrow band gap of ∼0.29 eV.^5–7 It has two Cu sites with different Cu–Se bond lengths and the Cu–Se framework provides paths for carrier transport with comparatively high mobility. Furthermore, a relatively complex crystal structure hinders the phonon transport effectively, resulting in a comparatively low thermal conductivity. So far, many efforts have been devoted to enhancing the TE performance of Cu₃SbSe₄, such as optimizing the carrier concentration or enhancing the Seebeck coefficient through doping using solid state reaction or chemical methods.^8–10 Tian Zhou et al. carried out research on 3D flower-like hierarchical Ti-doped Cu₃SbSe₄ microspheres, which possessed an ultra-low thermal conductivity of 0.38 W m⁻¹ K⁻¹ at 623 K and eventually reached the promising ZT value of ∼0.59 at this temperature.¹¹ Boyi Wang et al. developed a synergistic modulation of the thermal conductivity and power factor of Cu₃SbSe₄ based materials through Sn and Zr or Hf co-doping by using a facile microwave-assisted solvothermal method, yielding a peak ZT value of ∼0.82 for Cu₃Sb_0.91Sn_0.03Hf_0.06Se₄ at 623 K, which was ∼228% higher than that of pristine Cu₃SbSe₄.¹² Zhang Dan et al. carried out research on weakening p–d hybridization via Ag substitution at the Cu site, wherein the impaired Cu–Se bonding increased the hole concentration and the second phase AgSbSe₂ gave rise to large strain fluctuations, which reduced the lattice thermal conductivity remarkably. The peak ZT value of ∼ 0.9 was achieved in the sample Cu_2.85Ag_0.15SbSe₄. Xie Dandan et al. investigated (Ag and Sn) co-doped Cu₃SbSe₄ nanocrystals using the microwave-assisted solvothermal method. The results demonstrated that Cu_2−xSe nanoinclusions served as additional scattering centers for phonons. A maximum ZT of 1.18 at 623 K for Cu_2.8Ag_0.2Sb_0.95Sn_0.05Se₄ was obtained. We fabricated large-scale nanostructured Sn-doped samples through a low-temperature co-precipitation route and the highest ZT obtained for Cu₃Sb_0.98Sn_0.02Se₄ was 1.05.¹³ As an effective means to enhance the thermoelectric performance, low thermal conductivity was derived from increasing phonon scattering at grain boundaries and nanoparticles by nanostructuring.^14,15 In our previous work, second nanoparticles of Cu₃Sb_0.94Sn_0.06Se_2.5S_1.5 were introduced into the Cu₃Sb_0.96Sn_0.04Se₄ matrix and we obtained a minimum κ_L of about 0.62 W m⁻¹ K⁻¹ for the nanocomposites, which was about a 34% reduction compared with that of the matrix.¹⁶ Unfortunately, the introduction of nanoparticles resulted in a decrease of power factor and weakened the thermoelectric performance. Up to now, few effective methods can simultaneously realize a significant enhancement of power factor and reduction of thermal conductivity for the Cu₃SbSe₄ based system.

In order to further modulate electrical and thermal transport properties, we chose nanoparticles embedded in the Cu₃Sb_0.96Sn_0.04Se₄ matrix, introducing scattering centers, through which the power factor can be enhanced and κ_L can be decreased simultaneously. In this work, AgSb_0.98Ge_0.02Se₂ has been selected as the second nanophase because of its intrinsic low thermal conductivity and high TE performance.^17,18 The experimental results demonstrate that the incorporation of AgSb_0.98Ge_0.02Se₂ nanoinclusions in Cu₃Sb_0.96Sn_0.04Se₄ results in an enhancement of power factor due to the enhanced carrier mobility. In addition, thermal conductivity is reduced due to the enhanced scattering, which is consistent with our calculation of lattice thermal conductivity. As a result, a 9.5% enhancement of power factor as well as a 57.3% reduction of lattice thermal conductivity is realized in the composite sample with y = 3 wt%. Consequently, a record high ZT of 1.23 (at 675 K) is achieved, which is larger than that of any reported Cu₃SbSe₄-based materials.

2. Experimental section

Polycrystals of Cu₃Sb_0.96Sn_0.04Se₄ were prepared using a melting method. Elementary Cu (99.99%), Sb (99.5%), Sn (99.9%) and Se (99.95%) were precisely weighed according to the stoichiometric proportion and the mixed powders were sealed in evacuated quartz ampoules under a vacuum of 4 × 10⁻³ Pa. The ampoules were placed in a muffle furnace, slowly heated to 1173 K, kept at this temperature for 12 h, and then cooled to 773 K with the rate fixed at 1 K min⁻¹. After this, the tubes were quenched and then annealed at 573 K for 48 h to obtain uniform and homogeneous ingots. AgSb_0.98Ge_0.02Se₂ ingots were also prepared using a fusion method; the component elements, Ag (99.99%), Sb (99.5%), Sn (99.9%) and Se (99.95%), were sealed into evacuated quartz tubes at a pressure of 4.5 × 10⁻³ Pa. The samples were first heated to 723 K and held for one hour, then heated up to 1123 K and isothermally kept for 5 days, and finally cooled to room temperature. The resulting ingots were pulverized into a fine powder. Then, the obtained Cu₃Sb_0.96Sn_0.04Se₄ and AgSb_0.98Ge_0.02Se₂ powders were mixed on the basis of weight ratios of 1%, 3%, 5% and 7% by hand-grinding for 0.5 h. Finally, the bulk samples were obtained via spark plasma sintering (SPS) under a pressure of 50 MPa at 623 K for 5 minutes.

The room temperature phase constitution of the prepared powder was detected via X-ray diffraction using a Philips X’Pert PRO X-ray diffractometer with Cu Kα radiation (λ = 1.54056 Å). The operating voltage and current were kept at 40 kV and 400 mA, respectively. The morphology and fractographs were observed via field emission scanning electron microscopy (FE-SEM). In addition, the microstructure was investigated by using high-resolution transmission electron microscopy (HRTEM; JEOL JEM-2010) operating at a 200 kV accelerating voltage. The electrical properties and Seebeck coefficient were measured using the standard four-probe method with a ZEM-3 (ULVAC-RIKO) in a He atmosphere. Thermal conductivity κ was calculated by using the equation: κ = DdC_p, where D is the thermal diffusivity measured using a Netzsch LFA-457 in the temperature range from 300 K to 673 K. The density of the bulk sample (d) was measured using Archimedes’ method. The d values obtained for Cu₃Sb_0.96Sn_0.04Se₄–yAgSb_0.98Ge_0.02Se₂ (y = 0, 1, 3, 5, and 7 wt%) samples were 5.45 g cm⁻³, 5.55 g cm⁻³, 5.60 g cm⁻³, 5.51 g cm⁻³ and 5.61g cm⁻³. Special heat capacity Cp was determined via differential scanning calorimetry (DSC, PerkinElmer). Moreover, the Hall coefficient (R_H) was measured using the van der Pauw method under a magnetic field of 1.5T. The carrier concentration (p) and hall mobility (μ) of charge carriers were calculated from the relations p = 1/(R_He) and μ = R_H/p, where e is the electron charge. LabVIEW software was used to calculate the lattice thermal conductivity using the Callaway equation.

3. Result and discussion

The XRD patterns of AgSb_0.98Ge_0.02Se₂ and Cu₃Sb_0.96Sn_0.04Se₄-y AgSb_0.98Ge_0.02Se₂ (y = 0, 1, 3, 5, and 7 wt%) samples are shown in Fig. 1. One can see from the curves that all the diffraction peaks can be indexed well to the tetragonal structure phase of Cu₃Sb_0.96Sn_0.04Se₄ (JCPDS NO. 85-0003, space group I4̄2m) with lattice constants a = b = 5.63 Å and c = 11.23 Å. Moreover, in the patterns of the composite samples shown in curves from y = 3 wt% to y = 7 wt%, other than the reflection peaks from Cu₃Sb_0.96Sn_0.04Se₄, there are some small peaks at 2θ ≈ 31° and 44°, which become more evident with increasing AgSb_0.98Ge_0.02Se₂ content, and these peaks correspond to the (2 0 0) and (2 2 0) planes of AgSbSe₂ (JCPDS No. 12-0379, space group Fm3̄m) with lattice constants a = b = c = 5.786 Å, indicating that there are two phases in the composite samples.

Fig. 1 XRD patterns of the Cu₃Sb_0.96Sn_0.04Se₄ matrix and Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Ge_0.02Se₂ (y = 1, 3, 5, 7 wt%) composite samples. The inset shows the XRD pattern of AgSb_0.98Sn_0.02Se₂.

Fig. 2 shows the FESEM image of the fracture surface for the sample Cu₃Sb_0.96Sn_0.04Se₄-5 wt% AgSb_0.98Ge_0.02Se₂. One can observe from Fig. 2(a) that many nanoparticles (marked by yellow ellipses) with a size of about 50–100 nm are embedded well in the Cu₃Sb_0.96Sn_0.04Se₄ matrix, which agrees well with the EDS analysis and elemental mapping, where the bright zones correspond to the AgSb_0.98Ge_0.02Se₂-rich area. The SEM results confirmed that these particles are dispersed AgSb_0.98Ge_0.02Se₂, indicating that Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Ge_0.02Se₂ (y = 1, 3, 5, 7 wt%) composites were prepared well. In order to further explore the microstructure of the composite samples, we conducted TEM analysis on Cu₃Sb_0.96Sn_0.04Se₄–5 wt% AgSb_0.98Ge_0.02Se₂. As shown in Fig. 3(a), we observed that the particles with a size of around 200 nm are well embedded in the matrix with bright contrast. A high-resolution image is shown in Fig. 3(b).

Fig. 2 (a) Morphology of the bulk fracture of the sample Cu₃Sb_0.96Sn_0.04Se₄-5 wt% AgSb_0.98Ge_0.02Se₂; (b) EDS analysis of the areas marked in (a) by yellow ellipses; (c–h) EDX elemental mapping analysis of the sample Cu₃Sb_0.96Sn_0.04Se₄-5 wt% AgSb_0.98Ge_0.02Se₂ corresponding to (a).

Fig. 3 (a and b) Low magnification TEM image and HRTEM image for the bulk fracture of the sample Cu₃Sb_0.96Sn_0.04Se₄–5wt% AgSb_0.98Ge_0.02Se₂; (c and d) low magnification TEM image and HRTEM image of the Cu₃Sb_0.96Sn_0.04Se₄ matrix and dislocation distribution; the inset shows the SAED image of the dislocation areas.

The matrices with interplanar distances of 3.13 Å and 2.86 Å correspond to the lattice planes of (1 0 3) and (2 0 0) of the doped tetragonal Cu₃SbSe₄ compound. The lattice spacing of about 2.03 Å corresponds to the (2 2 0) plane of cubic AgSbSe₂. Obvious grain boundaries (GBs) are observed between the Cu₃SbSe₄ matrix and AgSbSe₂ inclusions. We also find the existence of a large number of dislocations and faults in the matrix (marked with arrows). Fig. 3(c) and (d) present the low resolution and HRTEM images of a dislocation, where we can see clear dislocation lines and crystal defects such as twin crystals. From the SAED pattern of the dislocation line shown in the inset of Fig. 3(d), we can observe that many light diffraction circles are regularly arrayed between the diffraction spots of tetragonal Cu₃SbSe₄, which indicates the distinguished feature of dislocations in the crystal.

The temperature dependence of electrical resistivity (ρ) for samples of Cu₃Sb_0.96Sn_0.04Se₄–yAgSb_0.98Ge_0.02Se₂ (y = 0, 1, 3, 5, and 7 wt%) is shown in Fig. 4(a). For all samples, ρ increases with increasing temperature, indicating representative degenerate semiconductor behavior. In addition, the existence of AgSb_0.98Ge_0.02Se₂ inclusions causes a remarkable decrease of ρ for samples of Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Ge_0.02Se₂ (y = 1, 3, 5, and 7 wt%) in the whole temperature range investigated. For example, the ρ value decreases from 2.8 × 10⁻⁵ Ω m to 2.05 × 10⁻⁵ Ω m with y increasing from 0 to 1 wt% at 300 K, respectively. However, with the increase of the nanoparticle content, ρ of the composite samples increases slightly and there is no significant change with the variation of the second phase content. The variation in carrier concentration (p) for samples of Cu₃Sb_0.96Sn_0.04Se₄–yAgSb_0.98Ge_0.02Se₂ (y = 0, 1, 3, 5, and 7 wt%) with temperature is shown in Fig. 5(a). It can be seen that the p value for all samples increases with increasing temperature, which is due to the thermal excitation of carriers. In addition, the p value decreases remarkably with a AgSb_0.98Ge_0.02Se₂ content of 1 wt%, while it increases with further increasing the AgSb_0.98Ge_0.02Se₂ content. For example, the p values at 300 K for samples with y = 0, 1, 3, 5, and 7 wt% are 1.32 × 10²⁰ cm⁻³, 6.59 × 10¹⁹ cm⁻³, 1.13 × 10²⁰ cm⁻³, 8.72 × 10¹⁹ cm⁻³, and 1.01 × 10²⁰ cm⁻³, respectively. When the AgSb_0.98Ge_0.02Se₂ nanoparticles are imbedded in the Cu₃Sb_0.96Sn_0.04Se₄ matrix, there are a lot of interfaces. A deep acceptor energy level will be formed, ascribed to interfacial defects, which can lead to a decrease in carrier concentration. However, due to the high carrier concentration of the AgSbSe₂ system as compared to that of the Cu₃SbSe₄ system,¹⁹ the higher content of AgSb_0.98Ge_0.02Se₂ embedded into the Cu₃Sb_0.96Sn_0.04Se₄ matrix will cause an increase of p as compared to the sample with y = 1 wt%.

Fig. 4 Temperature dependence of electrical resistivity, Seebeck coefficient and power factor for composite samples Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Sn_0.02Se₂ (y = 1, 3, 5, and 7 wt%) in the temperature range from 300 K to 675 K.

Fig. 5 (a) Temperature dependence of the carrier concentration; (b) plot of carrier mobility with temperature; (c and d) carrier concentration dependence of Seebeck coefficient (Pisarenko relation) and power factor at 300 K for all samples in this research.

The calculated values of carrier mobility (μ) are shown in Fig. 5(b). μ values for all composite samples are larger than that for the Cu₃Sb_0.96Sn_0.04Se₄ matrix. For instance, μ values at 300 K are 19 cm² (V s)⁻¹, 29.3 cm² (V s)⁻¹, 22.8 cm² (V s)⁻¹, 22.5 cm² (V s)⁻¹ and 24.8 cm² (V s)⁻¹ for samples with y = 0, 1, 3, 5 and 7 wt%, respectively. Meanwhile, we analyzed the relationship between ln μ and ln T. As shown in Fig. 5(c), the calculated slope (δ) ranges from −1.25 to −1.5, indicating that the main mechanism of carrier scattering results from lattice vibration. The decrease of carrier concentration due to a deep acceptor energy level causes an increase in mobility. As the AgSb_0.98Ge_0.02Se₂ content y is further increased, the decrease of mobility for the composite samples as compared to the sample with y = 1 wt% originates from the increasing grain boundaries. In this work, the trend of ρ with AgSb_0.98Sn_0.02Se₂ content is determined by the mobility μ values for samples Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Ge_0.02Se₂ (y = 0, 1, 3, 5, and 7 wt%).

The temperature dependence of S is shown in Fig. 4(b), and the Seebeck coefficients for all the measured samples are positive in the whole temperature range of 300–675 K, indicating that they are p-type semiconductors. The S values of all samples increase monotonically and S decreases obviously with the increase of nanoinclusion content from y = 1 to 3, 5, and 7 wt%. In order to understand the electrical transport properties better, the Mott formula is given as follows:²⁰

where k_B, E_f, and p are the Boltzmann constant, Fermi energy, and carrier concentration, respectively. From eqn (1), it can be seen that S is inversely proportional to carrier concentration and mobility. For the specimen Cu₃Sb_0.96Sn_0.04Se₄–1 wt% AgSb_0.98Ge_0.02Se₂, the carrier concentration is decreased, while the mobility μ increases largely, which causes a decrease of S value compared with the matrix. With increasing AgSb_0.98Ge_0.02Se₂ content y, the decrease of p determines the S value. The power factor calculated using the formula PF = S²/ρ is presented in Fig. 4(c). We can observe that the power factors for composite samples have greatly increased compared with Cu₃Sb_0.96Sn_0.04Se₄. In particular, PF for the sample Cu₃Sb_0.96Sn_0.04Se₄–5 wt% AgSb_0.98Ge_0.02Se₂ increases from 8.56 × 10⁻⁴ W m⁻¹ K⁻² to 1.3 × 10⁻³ W m⁻¹ K⁻² with increasing temperature from 300 K to 675 K, which is a 15% improvement as compared to the Cu₃Sb_0.96Sn_0.04Se₄ matrix in the whole temperature range measured. The largest power factor obtained in this work is 1.38 × 10⁻³ W m⁻¹ K⁻² for the sample Cu₃Sb_0.96Sn_0.04Se₄–5 wt% AgSb_0.98Ge_0.02Se₂, which is attributed to the increased mobility. The relationship between the carrier concentration and power factor is shown in Fig. 5(d). The result indicates that although carrier concentration decreases as compared to the matrix, the increase of carrier mobility still plays a leading role in enhancing the power factor.

Fig. 6(a) shows the temperature dependence of total thermal conductivity (κ_tot) for samples Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Ge_0.02Se₂ (y = 0, 1, 3, 5, and 7 wt%). κ_tot decreases with increasing temperature. Moreover, κ_tot decreases monotonically with increasing AgSb_0.98Ge_0.02Se₂ content (except for the sample with y = 7 wt%). For instance, κ_tot at 300 K drops from 3.98 W m⁻¹ K⁻¹ to 2.76, 2.33 and 2.02 W m⁻¹ K⁻¹ as y increases from 0 to 1, 3, and 5 wt%, respectively. In addition, as compared to the matrix, κ_tot of all composites decreases quickly at high temperatures (above 625 K). The results of the specific heat C_p measured via DSC for samples Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Ge_0.02Se₂ (y = 1, 3, 5, and 7 wt%) from 300 K to 675 K are shown in Fig. 6(b). One can see that the values of C_p for all samples are in the range from 0.32 to 0.37 J g⁻¹ K⁻¹, which is consistent with the theoretical value.²¹ To further analyze the thermal properties of the samples, the total thermal conductivity can be separated into two parts: carrier thermal conductivity and lattice thermal conductivity, as follows: κ_tot = κ_e + κ_L, where κ_e can be estimated using the Wiedemann–Franz relation, which is given by κ_e = LT/ρ, where the Lorenz number L is obtained using the Fermi integral function:²¹

In the SPB model, the scattering factor can be assumed as r = −1/2. So, the reduced Fermi energy (η) used above can be calculated from the experimental S values by using the following relationships:On the basis of the above calculation, the obtained L value ranges from 1.79 × 10⁻⁸ V² K⁻² to 1.86 × 10⁻⁸ V² K⁻², as shown in Fig. S1 (ESI†). As demonstrated in Fig. 6(c), κ_e initially increases with temperature and then slightly declines above 575 K for all samples. κ_L for Cu₃Sb_0.96Sn_0.04Se₄ follows a law of T⁻¹ (shown in Fig. 7(a)), indicating that phonon–phonon scattering (Umklapp scattering) dominates²² and κ_L for samples Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Ge_0.02Se₂ (y = 1, 3, 5, and 7 wt%) decreases across the whole temperature range investigated with increasing AgSb_0.98Ge_0.02Se₂ content y. Remarkably, κ_L for the sample with y = 5 wt% at 300 K is 1.79 W m⁻¹ K⁻¹, which is 53% smaller than that of the matrix sample. Nevertheless, κ_L for the sample with y = 7 wt% does not further decrease as y increases, which can be ascribed to the agglomeration of nanoparticles. To understand the reduction of the lattice thermal conductivity, we introduce the Debye model and the lattice thermal conductivity κ_L can be expressed using the Callaway equation aswhere h, θ, v, ζ, and τ are the Planck constant, Debye temperature, phonon velocity, usual dimensionless variable and phonon relaxation time, respectively. The phonon relaxation time is mainly related to scattering from multiple scattering centers in composite samples according to Matthiessen's rule, where τ can be written as²³

τ⁻¹ = τ_PD⁻¹ + τ_U⁻¹ + τ_N⁻¹ + τ_B⁻¹ + τ_D⁻¹ (7)

where τ_PD⁻¹, τ_U⁻¹, τ_N⁻¹, τ_B⁻¹ and τ_D⁻¹ are the relaxation times corresponding to scattering from point defects, phonon–phonon interactions, nanoparticles, boundaries and dislocations, respectively. For the Cu₃Sb_0.96Sn_0.04Se₄ matrix, the Umklapp process and point defects are the dominant mechanisms of phonon scattering. In the case of the nanocomposite samples, there are numerous nanoparticles playing a vital role in scattering low frequency phonons at low temperatures, which can explain the reason for the sharp decrease of κ_L near room temperature. Therefore, the combined scattering rate τ⁻¹ for samples Cu₃Sb_0.96Sn_0.04Se₄-y wt% AgSb_0.98Ge_0.02Se₂ (y = 1, 3, 5, and 7 wt%) can be written as

τ⁻¹ = τ_PD⁻¹ + τ_U⁻¹ + τ_B⁻¹ = Aω⁴ + Bω²T exp(−θ/3T) + ν/L (8)

where A and B are simplified parameters and L is the distance between nanoparticles (detailed calculation parameters are listed in Table S1, ESI†). As can be seen from the fitting results, with the increase of the content of the second phase, the phonon scattering from the point defects and nanoparticles increases first and then weakens, while the distance between nanoparticles initially decreases and then increases as the AgSb_0.98Ge_0.02Se₂ content y is increased to 7 wt%, which should be ascribed to the agglomeration of the second phase. As shown in Fig. 7(a), the calculated κ_L values using the U + PD + B model match well with the experimental value of κ_L for samples Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Ge_0.02Se₂ (y = 1, 3, 5, and 7 wt%). The results confirm that the phonon scattering from the grain boundaries of AgSb_0.98Ge_0.02Se₂ nanoparticles is the main reason for the reduction of κ_L compared with the matrix. The minimum lattice thermal conductivity observed in this work is consistent with the calculated limit of lattice thermal conductivity κ_min (around 0.4 W m⁻¹ K⁻¹),²⁴ which is 62% lower than that of the Cu₃Sb_0.96Sn_0.04Se₄ matrix.

Fig. 6 Temperature dependence of (a) total thermal conductivity (κ); (b) special heat capacity (C_p); (c) thermal conductivity contributed from carriers (κ_e); (d) lattice thermal conductivity (κ_L) for Cu₃Sb_0.96Sn_0.04Se₄ and Cu₃Sb_0.96Sn_0.04Se₄–yAgSb_0.98Ge_0.02Se₂ composite samples in the temperature range of 300–673 K.

Fig. 7 (a) Comparison between the experimental data and calculated values of lattice thermal conductivity; (b) ZT values for all samples in the whole temperature range; comparison of the reported (c) ZT_max and (d) ZT_ave values of Cu₃SbSe₄ based materials that were synthesized using a melting method or solid state reaction.

The ZT values for samples Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Ge_0.02Se₂ (y = 0, 1, 3, 5, and 7 wt%) are shown in Fig. 7(b), from which one can see that the temperature behavior of ZT for all five samples is similar: it increases with temperature. The ZT values of the composite samples increase greatly compared with that of the matrix. In particular, a maximum ZT of 1.23 is achieved at 675 K for the composite sample with y = 5 wt%, which is almost twice that of Cu₃Sb_0.96Sn_0.04Se₄. Clearly, this significant improvement of ZT can be attributed to the simultaneous enhancement of power factor and remarkable reduction of thermal conductivity. Among all the reported Cu₃SbSe₄ materials, our sample Cu₃Sb_0.96Sn_0.04Se₄–5 wt% AgSb_0.98Ge_0.02Se₂ exhibits the greatest thermoelectric performance (shown in Fig. 7(c)).^8,16,24–34 Meanwhile, we calculated the average ZT (ZT_ave) value of our sample and other samples according to the relationship:

where T_h and T_c are the temperatures of the hot-side and cold-side, respectively. As demonstrated in Fig. 7(d), the largest ZT_ave reaches 0.5 for the sample Cu₃Sb_0.96Sn_0.04Se₄-5 wt% AgSb_0.98Ge_0.02Se₂, which is 38% larger that of a previously reported composite sample.¹⁶

Conclusion

In this work, high thermoelectric performance was obtained in composite samples Cu₃Sb_0.96Sn_0.04Se₄-yAgSb_0.98Ge_0.02Se₂ (y = 1, 3, 5, and 7 wt%) obtained using melting and spark plasma sintering methods. The highest ZT of 1.23 was achieved for the sample Cu₃Sb_0.96Sn_0.04Se₄–3 wt% AgSb_0.98Ge_0.02Se₂ at 675 K. This outstanding improvement is attributed to the enhancement of power factor and the reduced lattice thermal conductivity due to the incorporation of AgSb_0.98Ge_0.02Se₂ inclusions. The obtained ZT = 1.23 is almost twice that of the Cu₃Sb_0.96Sn_0.04Se₄ matrix, which demonstrates that imbedding the second phase AgSb_0.98Ge_0.02Se₂ is a promising strategy to enhance the thermoelectric properties of the Cu₃SbSe₄ system.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work was supported by funding from the Natural Science Foundation of China under grant No. 51672278 and 11674322. We acknowledge Dr Z.W. Chen and Prof. Y. Z. Pei of Tong Ji University for their help in the measurement of Hall coefficient.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/d0qm00553c

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