KR20220068078A - Tetrahedrite-based thermoelectric materials and method for preparing the same - Google Patents

Abstract

The present invention provides a thermoelectric material with improved thermoelectric performance and a method of manufacturing the same. A thermoelectric material according to various embodiments of the present invention is a thermoelectric material in which Cu_(12)Sb_4S_(13) is doped with at least one selected from the group consisting of Ge, Si and Sn. In more detail, the thermoelectric material according to the present invention is Cu_(12)Sb_(4-y)M_yS_(13), where 0 < y <= 1.5, and M is at least one selected from the group consisting of Ge, Si and Sn. A method of manufacturing a thermoelectric material according to various embodiments of the present invention includes the steps of: preparing raw materials by weighing them according to the composition formula Cu_(12)Sb_(4-y)M_yS_(13) (where 0 < y <= 1.5, and M is at least one selected from the group consisting of Ge, Si and Sn); synthesizing a powder by mechanically alloying the prepared raw materials; and hot pressing (HP) the powder.

Description

Tetrahedrite-based thermoelectric materials and method for preparing the same

The present invention relates to a tetrahedrite-based thermoelectric material and a method for manufacturing the same, and more particularly, to a tetrahedrite-based thermoelectric material doped with Ge, Si or Sn, and a method for manufacturing the same.

Thermoelectric conversion technology that can directly convert thermal energy and electric energy is emerging as a promising eco-friendly energy system. Thermoelectric materials have different temperature ranges that show maximum performance, and it is known that Bi-Te-based materials in the room temperature range, Pb-Te-based, skutterudite-based materials in the medium temperature range, and Si-Ge-based materials in the high temperature range are known to be excellent. . However, there is a problem that most of them consist of toxic heavy metals or rare elements with limited reserves. Accordingly, a tetrahedrite compound composed of Cu and S, which has abundant reserves, is inexpensive, and is relatively light and nontoxic, is attracting attention as a promising p-type thermoelectric material in the mesophilic region. The main reason why tetrahedrite is attracting attention is because it shows low thermal conductivity and has high energy band degeneracy due to its symmetrical crystal structure, so it is useful for improving the output factor.

To obtain a high ZT according to the thermoelectric figure of merit relation, ZT = α ² σκ ^-1 T (α: Seebeck coefficient, σ: electrical conductivity, κ: thermal conductivity, T: absolute temperature), a high output factor (α ² σ) and Low thermal conductivity is required. However, since the Seebeck coefficient and the electrical conductivity have a trade-off relationship depending on the carrier concentration, the output factor is generally improved by optimizing the carrier concentration by doping. It is known that undoped tetrahedrite has a high hole concentration because the Fermi level is located in the valence band. Therefore, many research results have been conducted on the suppression of electron thermal conductivity and optimization of output factors by reducing the hole concentration by substituting a divalent or trivalent transition metal at the Cu ⁺ site. Heo et al. obtained ZT = 1.13 @ 575 K from Mn-substituted Cu ₁₁ MnSb ₄ S ₁₃ with a combination of high output factor and low thermal conductivity, and Lu et al. obtained ZT = 1 @ 720 K in Cu ₁₁ ZnSb ₄ S ₁₃ . values have been announced.

An object of the present invention is to provide a thermoelectric material with improved thermoelectric performance and a method for manufacturing the same by doping Ge, Si or Sn to reduce thermal conductivity and optimize output factors by doping tetrahedrite.

A thermoelectric material according to various embodiments of the present disclosure is a thermoelectric material in which Cu ₁₂ Sb ₄ S ₁₃ is doped with at least one selected from the group consisting of Ge, Si, and Sn. More specifically, the thermoelectric material of the present invention is Cu ₁₂ Sb _{4 -y} M _y S ₁₃ , where 0 < _y ≤ 1.5, and M is at least one selected from the group consisting of Ge, Si, and Sn.

In the method of manufacturing a thermoelectric material according to various embodiments of the present invention, the raw material is prepared by the composition formula Cu ₁₂ Sb _{4 - y} M _y S ₁₃ (where 0 < y ≤ 1.5, and M is from the group consisting of Ge, Si and Sn Preparing by weighing according to the selected at least one); synthesizing powder by mechanically alloying the prepared raw material; and hot pressing (HP) the powder.

Various embodiments of the present invention may provide a thermoelectric material with improved thermoelectric performance and a method for manufacturing the same.

According to an embodiment of the present invention, Ge is doped tetrahedrite. In the case of Cu ₁₂ Sb _3.8 Ge _0.2 S ₁₃ , a maximum value of ZT = 0.74 can be obtained at 723 K, and Cu ₁₂ Sb ₃ , a Si-doped tetrahedrite _{. 8} Si _{0 . 2} For S ₁₃ At 723 K, a maximum value of ZT = 0.63 can be obtained, and Sn-doped tetrahedrite Cu ₁₂ Sb _{3 . 9} Sn _{0 . 1} S ₁₃ In the case of , a maximum value of ZT = 0.66 can be obtained at 723 K.

The manufacturing method of the present invention can prevent volatilization of elemental components and segregation of dopants that may occur in the dissolution process for the conventional tetrahedrite-based compound production by using mechanical alloying and hot compression molding, and a homogeneous tetrahedrite phase can be synthesized. In addition, it is very suitable for mass production by lowering the manufacturing cost and remarkably reducing the process time to increase the process efficiency compared to the dissolution method in which a high temperature must be maintained for a long time. In addition, it is an effective process because it can _produce homogeneous Cu ₁₂ Sb _{4 -y} M _y S ₁₃ through solid-state synthesis within a relatively short time.

1 shows the X-ray diffraction pattern of Example 1.
FIG. 2 is an SEM photograph of the fracture surface of Example 1. FIG.
Figure 3 is Cu ₁₂ Sb _{3 . 8} Ge _{0 . 2} This is a picture of EDS element mapping of the S ₁₃ specimen.
4 is a graph of the Seebeck coefficient of Example 1.
5 is a graph of electrical conductivity of Example 1.
6 shows an output factor (PF = α ² σ ) of Example 1.
7 shows the thermal conductivity of Example 1.
8 shows the Lorenz number of Example 1.
9 shows the dimensionless thermoelectric figure of merit ( ZT ) of Example 1.
10 shows the X-ray diffraction pattern of Example 2.
11 is an SEM photograph of the fracture surface of Example 2. FIG.
12 is Cu ₁₂ Sb _{3 . 8} Si _{0 . 2} This is a picture of EDS element mapping of the S ₁₃ specimen.
13 is a graph of the Seebeck coefficient of Example 2.
14 is a graph of electrical conductivity of Example 2.
15 shows an output factor (PF = α ² σ ) of Example 2.
16 shows the thermal conductivity of Example 2.
17 shows the Lorenz number of Example 2.
18 shows the dimensionless thermoelectric figure of merit ( ZT ) of Example 2.
19 shows the X-ray diffraction pattern of Example 3.
20 is an SEM photograph of the fracture surface of Example 3.
21 is Cu ₁₂ Sb _{3 . 9} Sn _{0 . 1} This is an EDS element mapping picture of the S ₁₃ specimen.
22 is an electrical conductivity graph of Example 3. FIG.
23 is a graph of the Seebeck coefficient of Example 3.
24 shows an output factor (PF = α ² σ ) of Example 3.
25 shows the thermal conductivity of Example 3.
26 shows the Lorenz number of Example 3.
27 shows the dimensionless thermoelectric figure of merit ( ZT ) of Example 3.

Hereinafter, various embodiments of the present document will be described with reference to the accompanying drawings. The examples and terms used therein are not intended to limit the technology described in this document to a specific embodiment, but it should be understood to cover various modifications, equivalents, and/or substitutions of the embodiments.

A thermoelectric material according to various embodiments of the present disclosure is a thermoelectric material in which Cu ₁₂ Sb ₄ S ₁₃ is doped with at least one selected from the group consisting of Ge, Si, and Sn. More specifically, the thermoelectric material of the present invention is Cu ₁₂ Sb _{4 -y} M _y S ₁₃ , where 0 < _y ≤ 1.5, and M is at least one selected from the group consisting of Ge, Si, and Sn.

The thermoelectric material according to various embodiments of the present invention is Cu ₁₂ Sb _{4 -y} Ge _y S ₁₃ , where _y is 0.1, 0.2, 0.3 or 0.4.

The thermoelectric material according to various embodiments of the present invention is Cu _{12 Sb 4 -y} Si _y S ₁₃ , where y is 0.1, 0.2, 0.3 or 0.4.

The thermoelectric material according to various embodiments of the present invention is Cu ₁₂ Sb _{4 -y} Sn _y S ₁₃ , where _y is 0.1, 0.2, 0.3 or 0.4.

According to various embodiments of the present disclosure, it is possible to provide a thermoelectric material with improved thermoelectric performance by reducing thermal conductivity and optimizing an output factor.

A method of manufacturing a thermoelectric material according to various embodiments of the present disclosure includes preparing a raw material; synthesizing powder by mechanically alloying the prepared raw material; and hot pressing (HP) the powder.

In the step of preparing the raw material, the raw material is prepared according to the compositional formula Cu ₁₂ Sb _{4 - y} M _y S ₁₃ (where 0 < y ≤ 1.5, and M is at least one selected from the group consisting of Ge, Si and Sn) It can be weighed and prepared. At this time, Cu may prepare an elemental powder having a particle size of less than 45 μm, Sb may prepare an elemental powder having a particle size of less than 150 μm, and S may prepare an elemental powder having a particle size of less than 75 μm. In addition, Ge may prepare an elemental powder having a particle size of less than 45 μm, Si may prepare an elemental powder having a particle size of less than 75 μm, and Sn may prepare an elemental powder having a particle size of less than 35 μm.

Next, in the step of synthesizing the powder, the prepared raw material may be mechanically alloyed (MA). The prepared mixed powder and a steel ball having a diameter of 4 to 6 mm can be mixed in a ratio of 1:15 to 1:20 for ball milling. Specifically, it may be ball milled at 200 rpm to 2000 rpm for 1 hour to 100 hours. Cu ₁₂ Sb _{4 - y} M _y S ₁₃ for ball milling speed less than 200 rpm The phase may not be synthesized properly, and when it exceeds 2000 rpm, phase decomposition may occur or a secondary phase may be formed due to excessive energy. On the other hand, if the ball milling time is shorter than 1 hour, the Cu ₁₂ Sb _{4 -y} M _y S ₁₃ phase _may not be properly synthesized, and if it is longer than 100 hours, phase decomposition may occur or a secondary phase may be formed due to excessive energy. . Preferably, it can be ball milled at 350 rpm for 22 hours to 26 hours.

Next, the hot compression molding is preferably performed in a temperature range of 473 K to 873 K and a pressure range of normal pressure to 100 MPa. At a condition lower than this range, a desired sintering result may not be obtained, and at a condition higher than this range, the manufacturing cost may increase. On the other hand, preferably, in the case of sintering by hot pressing (HP), it may be sintered at a pressure of 70 Mpa in a temperature range of 523 K to 623 K. More preferably, it may be sintered at a temperature of 573 K and a pressure of 70 Mpa.

The manufacturing method of the present invention can prevent volatilization of elemental components and segregation of dopants that may occur in the dissolution process for the conventional tetrahedrite-based compound production by using mechanical alloying and hot compression molding, and a homogeneous tetrahedrite phase can be synthesized. In addition, it is very suitable for mass production by lowering the manufacturing cost and remarkably reducing the process time to increase the process efficiency compared to the dissolution method in which a high temperature must be maintained for a long time. In addition, it is an effective process because it can _produce homogeneous Cu ₁₂ Sb _{4 -y} M _y S ₁₃ through solid-state synthesis within a relatively short time.

Hereinafter, various embodiments of the present invention will be described in detail with reference to the accompanying drawings.

Example 1: Tetrahedrite Cu Doped with Ge ₁₂ Sb _4-y Ge _y S ₁₃

Ge-doped tetrahedrite Cu ₁₂ Sb _{4 -y} Ge _y S ₁₃ ( _y = 0.1, 0.2, 0.3, and 0.4) was synthesized by mechanical alloying (MA). Cu (purity 99.9%, < 45 μm, Kojundo), Sb (purity 99.999%, < 150 μm, Kojundo), Ge (purity 99.99%, < 45 μm, Kojundo), S (purity 99.99%, < 75 μm, Kojundo) ) were accurately weighed according to the stoichiometric composition ratio, and the weight ratio (Ball to Powder Ratio, BPR) of the 5 mm diameter stainless steel ball and the mixed powder was set to 20 and put into the hardened steel jar. After making the inside of the jar in a vacuum state, Ar gas was injected, and mechanical alloying was performed for 24 hours using a planetary mill (Fritsch Pulverisette5) at a rotation speed of 350 rpm. The powder synthesized through MA was charged into a graphite mold with an inner diameter of 10 mm and sintered by HP at a temperature of 723 K and a pressure of 70 MPa in a vacuum atmosphere for 2 hours. The sintered specimen was subjected to a disk of 10 mm (diameter) Х 1 mm (thickness) to measure thermal conductivity and Hall modulus and 3 mm (horizontal) Х 3 mm (length) Х 9 mm ( height) by cutting and polishing the disc.

[Experimental Example 1]

Experimental method

)와 무차원 성능지수(ZT = α²σκ^-1T)를 평가하였다.Phase analysis and crystallinity of the sintered specimen were confirmed through X-ray diffraction analysis (XRD; Bruker D8-Advance) using Cu Kα radiation (λ = 0.15405 nm). A diffraction pattern was obtained by measuring at a diffraction angle of 2θ = 10-90° with a step size of 0.02° and a scan speed of 0.4 sec/step. The lattice constant was calculated using Rietveld refinement (TOPAS). The microstructure of the fracture surface of the sintered specimen was observed using a scanning electron microscope (SEM; FEI Quanta400) and EDS elemental energy levels were Cu-Lα (0.928 eV) and Sb, respectively, using an energy-dispersive spectrometer (EDS; Bruker Quantax200). -Lα (3.604 eV), Ge-Lα (1.188 eV), and S-Kα (2.309 eV) were analyzed and the constituent elements were analyzed. A homogeneous distribution of elements was confirmed through mapping. The Hall coefficient, carrier concentration and mobility of the sintered body were measured by applying a constant 1 T magnetic field and 100 mA current at room temperature using the van der Pauw (Keithley 7065) method. In the temperature (T) range of 323 to 723 K, Seebeck coefficient (α) and electrical conductivity (σ) were measured using the temperature differential method and DC 4-terminal method in He atmosphere, respectively (Ulvac-Riko ZEM-3), and the thermal conductivity (κ) is the output factor (PF =

) and dimensionless figure of merit (ZT = α ² σκ ^-1 T) were evaluated.

Experimental Example 1-1: X-ray diffraction pattern analysis of Example 1

1 shows an X-ray diffraction pattern according to the Ge content of Ge-doped _tetrahedrite Cu ₁₂ Sb _{4 -y} Ge _y S ₁₃ according to Example 1. FIG.

The peaks of all specimens were consistent with the standard diffraction data (ICDD PDF#024-1318), indicating that the tetrahedrite single phase was free of residual elements and secondary phases.

Meanwhile, the lattice constants calculated through Rietveld refinement are shown in Table 1 below.

Composition Relative density
[%] Lattice constant
[nm] Lorenz number
[10 ^-8 V ² K ^-2 ] Nominal Actual Cu ₁₂ Sb _3.9 Ge _0.1 S ₁₃ Cu _12.89 Sb _3.83 Ge _0.12 S _12.15 99.3 1.0343 1.74 Cu ₁₂ Sb _3.8 Ge _0.2 S ₁₃ Cu _12.42 Sb _3.72 Ge _0.20 S _12.64 100.0 1.0339 1.69 Cu ₁₂ Sb _3.7 Ge _0.3 S ₁₃ Cu _12.32 Sb _3.90 Ge _0.25 S _12.51 99.6 1.0337 1.67 Cu ₁₂ Sb _3.6 Ge _0.4 S ₁₃ Cu _12.44 Sb _3.90 Ge _0.41 S _12.23 98.5 1.0334 1.61

Referring to Table 1, as a result of Rietveld analysis, the lattice constant decreased from 1.0343 nm to 1.0334 nm as the Ge content increased. This means that the ionic radius (39 pm) of Ge ⁴⁺ is ^that of Sb ³⁺ Since it is smaller than the ionic radius (76 pm), it seems that the lattice constant decreases as the Ge content increases.

Experimental Example 1-2: Analysis of the shape and composition of Example 1

Figure 2 is a sintered Cu ₁₂ Sb _{4 - y} Ge _y S ₁₃ of the specimen This is an SEM picture of the fracture surface. The difference in microstructure according to the Ge content was not observed, and a densely sintered microstructure was exhibited with few pores. On the other hand, as shown in Table 1 above, all the specimens showed a high relative density of 98.5-100.0%, and the nominal composition and the actual composition were similar. . FIG. 3 shows a photograph of EDS element mapping of a Cu ₁₂ Sb _3.8 Ge _0.2 S ₁₃ specimen, confirming that all constituent elements were homogeneously distributed.

Experimental Example 1-3: Seebeck coefficient analysis of Example 1

= (8/3)ð ² k _B ² m ^* Te ^-1 h ^{- 2}(ð/3n)^2/3 (k_B: Boltzmann constant, m^*: effective carrier mass, e: electronic charge, h: Planck constant, n: carrier concentration)로 나타내어진다. y = 0.1-0.3의 경우, 온도가 증가함에 따라 Seebeck 계수가 증가하였지만 y = 0.4의 경우 온도가 증가함에 따라 Seebeck 계수가 감소하였다. 온도가 증가하면 p-type 반도체의 Seebeck 계수 값이 양의 방향으로 상승하다가 특정 온도 이상이 되면 진성 천이가 발생하여 carrier 농도가 급격히 증가하면서 Seebeck 계수는 감소하게 된다. 즉, 반도체의 온도가 증가할수록 Seebeck 계수가 특정 온도에서 최고 값을 보인 후 감소하게 된다. 본 실험예에서는 Cu₁₂Sb_3.6Ge_0.4S₁₃를 제외한 나머지 시편은 측정 온도 범위 안에서 진성 전이가 발생하지 않았으므로 온도 증가에 따라 Seebeck계수가 증가하였지만 Cu₁₂Sb_{3 . 6}Ge_{0 . 4}S₁₃의 경우 323 K이하의 온도에서 미리 진성 천이가 발생하게 되어 온도 증가에 따라 carrier 농도가 급격히 증가하여 Seebeck 계수는 감소하였다. 일정한 온도에서 Ge의 함량이 증가할수록 Seebeck 계수가 증가하였는데, 이는 캐리어 농도와 Seebeck 계수는 반비례 관계이므로 Ge의 함량이 증가할수록 주된 캐리어인 홀을 감소시켜 Seebeck 계수가 증가하였다. 본 실험예에서 Cu₁₂Sb_{3 . 6}Ge_{0 . 4}S₁₃의 경우 323-723 K온도 범위에서 251-224 μVK^{- 1}으로 가장 높은 Seebeck 계수의 값을 나타내었다. 한편, 타 연구에 따르면 300-690 K의 온도 범위에서 도핑 하지 않은 Cu₁₂Sb₄S₁₃는 105-140 μVK^-1의 Seebeck 계수를 나타내었는데 본 실시예는 이보다 높은 값을 보였다. 또한, 타 연구에 따르면 Cu₁₂Sb_{4 - y}Bi_yS₁₃ (y = 0.1-0.4)의 Cu₁₂Sb_{3 . 9}Bi_{0 . 1}S₁₃이 323-723 K에서 최대 154-187 μVK^{- 1}를 나타내었고, Cu₁₂Sb_{2 . 15}Te_{1 . 85}S₁₃ 의 시편은 300-700 K에서 최대 190-260 μVK^-1을 나타낸 결과가 있어, Sb 자리에 Ge의 도핑은 테트라헤드라이트의 Seebeck 계수 증가에 매우 효과적임을 알 수 있다.4 _shows the temperature dependence of the Seebeck coefficient according to the Ge content of Cu ₁₂ Sb _{4 -y} Ge _y S ₁₃ . All specimens showed positive Seebeck coefficient values like Hall coefficients, indicating that the main carriers were p-type semiconductors with holes. Seebeck coefficient in p-type semiconductor is

= (8/3) ð ² k _B ² m ^* Te ^-1 h ^{- 2} ( ð /3 n ) ^2/3 (k _B : Boltzmann constant, m ^* : effective carrier mass, e: electronic charge, h: Planck constant, n: carrier concentration). In the case of y = 0.1-0.3, the Seebeck coefficient increased as the temperature increased, but in the case of y = 0.4, the Seebeck coefficient decreased as the temperature increased. When the temperature increases, the Seebeck coefficient value of the p-type semiconductor rises in the positive direction, but when it exceeds a certain temperature, an intrinsic transition occurs and the carrier concentration rapidly increases and the Seebeck coefficient decreases. That is, as the temperature of the semiconductor increases, the Seebeck coefficient shows the highest value at a specific temperature and then decreases. In this experimental example, except for Cu ₁₂ Sb _3.6 Ge _0.4 S ₁₃ , no intrinsic transition occurred within the measurement temperature range, so the Seebeck coefficient increased as the temperature increased, but Cu ₁₂ Sb _{3 . 6} Ge _{0 .} In the case of ₄ S ₁₃ , an intrinsic transition occurred in advance at a temperature below 323 K, and as the temperature increased, the carrier concentration rapidly increased and the Seebeck coefficient decreased. At a constant temperature, the Seebeck coefficient increased as the Ge content increased. This is because the carrier concentration and the Seebeck coefficient are in inverse proportion to each other. As the Ge content increased, the Seebeck coefficient increased by reducing the hole, the main carrier. In this experimental example, Cu ₁₂ Sb _{3 . 6} Ge _{0 . 4} In the case of S ₁₃ , it was 251-224 μVK ^{- 1} in the 323-723 K temperature range, showing the highest Seebeck coefficient value. On the other hand, according to another study, undoped Cu ₁₂ Sb ₄ S ₁₃ in the temperature range of 300-690 K showed a Seebeck coefficient of 105-140 μVK ^-1 , but this example showed a higher value than this. In addition, according to other studies, Cu ₁₂ Sb _{4 - y} Bi _y S ₁₃ Cu ₁₂ Sb ₃ in (y = 0.1-0.4) _{. 9} Bi _{0 . 1} S ₁₃ exhibited a maximum of ^154-187 μVK ^-1 at 323-723 K, and Cu ₁₂ Sb _{2 . 15} Te _{1 .} The specimen of ₈₅ S ₁₃ showed a maximum of 190-260 μVK ^-1 at 300-700 K, suggesting that doping of Ge at the Sb site is very effective in increasing the Seebeck coefficient of tetrahedrite.

Experimental Example 1-4: Electrical conductivity analysis of Example 1

5 _shows the temperature dependence of the electrical conductivity of Cu ₁₂ Sb _{4 -y} Ge _y S ₁₃ . In the case of y = 0.1-0.3 specimens, the electrical conductivity increased as the temperature increased, and then decreased slightly above 623 K. It is expected that the electrical conductivity of tetrahedrite is changed from non-degenerate semiconductor to metallic (degenerate semiconductor) according to temperature. In the case of the y = 0.4 specimen, the non-degenerate semiconductor behavior was maintained in the measurement temperature range. At a certain temperature, the electrical conductivity decreased as the Ge content increased. This is because the electrical conductivity is reduced due to the reduction of holes, which are the main carriers, due to electrons added by doping with Ge. In this experimental example, Cu ₁₂ Sb _{3 . 9} Ge _{0 . 1} S ₁₃ obtained the highest electrical conductivity at 323-723 K as (2.8-3.2) x 10 ⁴ Sm ^-1 .

Experimental Example 1-5: Analysis of output factors of Example 1

의 식으로 나타내어지며, Seebeck 계수와 전기전도도에 비례하여 증가한다. 캐리어 농도는 전기전도도와 Seebeck 계수에 서로 반대의 영향을 주기 때문에 출력인자 값을 최대로 얻기 위해서는 최적의 캐리어 농도를 찾는 것이 요구된다. 측정 온도 범위 안에서는 온도가 증가함에 따라 출력 인자는 증가하였다. 이는 제백계수(도 4)와 전기전도도(도 5)의 온도 의존성의 결과이다. 일정 온도에서 Ge의 함량이 증가함에 따라서 출력인자는 감소하였다, 이는 Ge 도핑에 의한 Seebeck 계수의 증가 보다는 전기전도도의 감소 효과가 더 우세하였기 때문이다. Cu₁₂Sb_3.9Ge_0.1S₁₃시편은 가장 낮은 Seebeck 계수 값을 가졌음에도 불구하고, 가장 높은 전기전도도 값의 영향으로 인해 323-723 K의 온도 범위에서 가장 높은 출력인자 값인 0.38-0.87 mWm^-1K^-2을 얻어냈다. 6 shows the temperature dependence of Cu ₁₂ Sb _{4 -y} Ge _y S ₁₃ on the output factor ₍ PF). The output factor is PF =

It is expressed by the formula, and it increases in proportion to the Seebeck coefficient and electrical conductivity. Since carrier concentration has opposite effects on electrical conductivity and Seebeck coefficient, it is required to find the optimal carrier concentration to obtain the maximum value of the output factor. Within the measurement temperature range, the output factor increased as the temperature increased. This is a result of the temperature dependence of the Seebeck coefficient (FIG. 4) and the electrical conductivity (FIG. 5). The output factor decreased as the content of Ge increased at a certain temperature, because the effect of decreasing the electrical conductivity was more dominant than the increase of the Seebeck coefficient by Ge doping. Although the Cu ₁₂ Sb _3.9 Ge _0.1 S ₁₃ specimen had the lowest Seebeck coefficient value, the highest output factor value in the temperature range of 323-723 K was 0.38-0.87 mWm ^-1 K due to the influence of the highest electrical conductivity value. ^-2 was obtained.

Experimental Example 1-6: Thermal conductivity analysis of Example 1

식을 사용하여 계산하였다. 도 8과 같이 323-723 K에서 온도가 증가함에 따라 y = 0.1-0.3의 경우 Lorenz number가 감소하였지만, y = 0.4의 경우 온도가 증가함에 따라 약간 증가하였다. 일정 온도에서는 Ge 함량이 증가함에 따라 Lorenz number가 감소하였다. Lorenz number의 감소는 비축퇴 반도체의 거동을 의미한다. Cu₁₂Sb_{3 . 6}Ge_{0 . 4}S₁₃ 시편의 경우 (1.61-1.64) x 10^-8 V²K^{- 2}으로 가장 낮은 값을 보였다. 7 ₍ a) is a graph showing the temperature dependence of the thermal conductivity of Cu ₁₂ Sb _{4 -y} Ge _y S ₁₃ . As the temperature increased, the thermal conductivity increased and then decreased after 623 K. When the Ge content increases at a constant temperature, the thermal conductivity decreases, Cu ₁₂ Sb _{3 . 6} Ge _{0 . 4} S ₁₃ showed a minimum value of 0.46-0.59 Wm ^-1 K ^-1 at 323-723 K. Thermal conductivity is the sum of electronic thermal conductivity (κ _E ) and lattice thermal conductivity (κ _L ), and electronic thermal conductivity was calculated according to the Wiedemann-Franz law (κ _E = LσT, L: Lorenz number). where Lorenz number [10 ^-8 V ² K ^{- 2} ] is

It was calculated using the formula. As shown in FIG. 8, as the temperature increased from 323-723 K, the Lorenz number decreased in the case of y = 0.1-0.3, but slightly increased as the temperature increased in the case of y = 0.4. At a constant temperature, the Lorenz number decreased as the Ge content increased. A decrease in the Lorenz number means the behavior of a non-degenerate semiconductor. Cu ₁₂ Sb _{3 . 6} Ge _{0 .} In the case of ₄ S ₁₃ specimen, (1.61-1.64) x 10 ^-8 V ² K ^{- 2} showed the lowest value.

As shown in FIG. 7(b), the electronic thermal conductivity increased as the temperature increased, but the total thermal conductivity of FIG. 7(a) decreased after showing a maximum value at 623 K. This is a result of the competition between increasing electronic thermal conductivity and decreasing lattice thermal conductivity with increasing temperature. Electronic thermal conductivity is related to electrical conductivity (carrier concentration) and Lorenz number, and it coincides with the tendency to decrease with increasing Ge content at a certain temperature, Cu ₁₂ Sb _{3 . 6} Ge _{0 . 4 In} the case of the S13 specimen, the lowest value of electronic thermal conductivity was shown at 323-723 K ^, 0.01-0.08 Wm ^-1 K ^-1 . In this experimental example, the lattice thermal conductivity decreased as the Ge content increased, and Cu ₁₂ Sb _{3 . 6} Ge _{0 . 4 In} the case of the S13 specimen, the lowest value was shown at 323-723 K at ^0.45-0.55 Wm ^-1 K ^-1 . For this reason, as shown in FIG. 7(a), the overall thermal conductivity showed the lowest value of 0.46-0.59 Wm ^-1 K ^-1 .

Experimental Example 1-7: Thermoelectric figure of merit (ZT) analysis of Example 1

Figure 9 is Cu ₁₂ Sb _{4 - y} Ge _y S ₁₃ fabricated by the MA-HP process according to Example 1 This is the result of comparing the ZT value of (y = 0.1-0.4) and the ZT value of the specimen doped with other elements (Te, Bi, and S) at the Sb site. For Cu ₁₂ Sb _{4 - y} Ge _y S ₁₃ , the ZT values of all specimens increased as the temperature increased. This is a result of the increase of the output factor according to the increase in temperature and the maintenance of low thermal conductivity. At a constant temperature, the high ZT value was maintained as the Ge doping content increased, but when y ≥ 0.3, the ZT value decreased. This is because the output factor decreased as the Ge content increased. Cu ₁₂ Sb ₃ , due to the high power factor (0.77 mWm ^-1 K ^{- 2} ) and low thermal conductivity (0.73 Wm ^-1 K ^{- 1} ) _{. 8} Ge _{0 . 2} S ₁₃ showed the highest ZT = 0.74 at 723 K. The synthesis using mechanical alloying in this example is a relatively short time of 24 hours, and since it can be synthesized in a single phase without additional heat treatment thereafter, it was possible to produce strong and excellent Ge-doped tetrahetrite in a short time with excellent thermoelectric performance.

Example 2: Si Doped Tetrahedrite Cu ₁₂ Sb _4-y Si _y S ₁₃

Si-doped Tetrahedrite Cu ₁₂ Sb _{4 - y} Si _y S ₁₃ (y = 0.1, 0.2, 0.3, and 0.4) was synthesized with MA. Powdered elements Cu (purity 99.9%, < 45 μm, Kojundo), Sb (purity 99.999%, < 150 μm Kojundo), Si (purity 99.99% < 75 μm, Kojundo), S (purity 99.99%, < 75 μm) , Kojundo) was weighed according to the stoichiometric ratio, and the weight ratio of the raw material powder to the balls was 20, and mechanical alloying was performed for 24 hours at a rotation speed of 350 rpm using a 5mm STS ball and a hardened steel jar using a planetary mill (Fritsch Pulverisette5). carried out. The obtained powder was charged into a graphite mold having an inner diameter of 10 mm, and hot pressing was performed at a temperature of 723 K and a pressure of 70 MPa for 2 hours.

[Experimental Example 2]

Experimental Example 2-1: X-ray diffraction pattern analysis of Example 2

10 is an X-ray diffraction analysis result of a Cu ₁₂ Sb _{4 -y} Si _y S ₁₃ sintered _body manufactured by the MA-HP method. The tetrahedrite phase was successfully synthesized, and no Cu-Sb-S or Cu-S-based secondary phase was produced. However, Si residual elements appeared in the specimens with Si content y ≥ 0.3.

Meanwhile, the lattice constants calculated through Rietveld refinement are shown in Table 2 below.

Composition Relative density
[%] Lattice constant
[nm] Lorenz number
[10 ^-8 V ² K ^-2 ] Nominal Actual Cu ₁₂ Sb _3.9 Si _0.1 S ₁₃ Cu _12.50 Sb _3.82 Si _0.10 S _12.57 100.0 1.0357 1.86 Cu ₁₂ Sb _3.8 Si _0.2 S ₁₃ Cu _13.05 Sb _3.78 Si _0.22 S _11.93 99.1 1.0337 1.85 Cu ₁₂ Sb _3.7 Si _0.3 S ₁₃ Cu _13.22 Sb _3.55 Si _0.32 S _11.89 98.9 1.0337 1.81 Cu ₁₂ Sb _3.6 Si _0.4 S ₁₃ Cu _12.55 Sb _3.59 Si _0.39 S _12.46 100.0 1.0336 1.81

Referring to Table 2, the lattice constant calculated by Rietveld refinement (TOPAS Program) decreased from 1.0357 to 1.0336 nm as the Si content increased. It seems that the lattice constant decreased due to partial substitution of small Si ^{4 +} (26 pm) in Sb ^{3 +} (76 pm) with a large ionic radius. In the case of y ≥ 0.3, there is little change in the lattice constant, and it is judged as a solid solution limit for the Sb site, judging from the Si residue.

Experimental Example 2-2: Analysis of the shape and composition of Example 2

11 is an SEM image of observing the fracture surface of the sintered specimen. The difference in microstructure was not observed according to the Si content, and it was confirmed that the microstructure was densely sintered with few pores. As shown in Table 2, a high relative density of 98.9-100.0% was shown, and the nominal composition and the actual composition were similar. 12 is Cu ₁₂ Sb _{3 . 8} Si _{0 . 2} This is an EDS element mapping image of the S ₁₃ specimen, and it was confirmed that all constituent elements were uniformly distributed. Therefore, in the MA-HP process used as a solid-phase synthesis process, volatilization of component elements was suppressed, and a homogeneous Tetrahedrite phase could be synthesized and molded.

Experimental Example 2-3: Seebeck coefficient analysis of Example 2

)를 나타낸다. 모든 시편에서 양의 Seebeck 계수 값이 나타났으며, 이는 주된 캐리어가 홀(hole)인 p-type 반도체임을 의미한다. 723 K까지 온도가 증가할 수록 Seebeck 계수는 증가하였고 Si이 도핑된 Tetrahedrite의 진성천이 온도는 723 K 보다 훨씬 높은 것으로 판단된다. 진성 천이가 발생하면 캐리어 농도가 급격히 증가하여 Seebeck 계수가 감소하기 때문이다. 일정 온도에서 Si 함량이 증가할수록 Seebeck 계수가 증가하였다. 323-623 K의 온도범위에서는 Cu₁₂Sb_{3 . 6}Si_{0 . 4}S₁₃의 시편이 136-169 μVK^{- 1} 의 Seebeck 계수 값을 얻었고 723K 에서는 Cu₁₂Sb_{3 . 7}Si_{0 . 3}S₁₃에서 178 μVK^-1으로 가장 높은 Seebeck 계수값을 나타내었다. Thermoelectric properties were measured in a temperature range of 323-723K. 13 is a Seebeck coefficient of Cu ₁₂ Sb _{4 -y} Si _y S _{13 (}

) is indicated. All specimens showed positive Seebeck coefficient values, indicating that the predominant carriers were p-type semiconductors with holes. As the temperature increased up to 723 K, the Seebeck coefficient increased. This is because, when an intrinsic transition occurs, the carrier concentration increases rapidly and the Seebeck coefficient decreases. As the Si content increased at a constant temperature, the Seebeck coefficient increased. In the temperature range of 323-623 K, Cu ₁₂ Sb _{3 . 6} Si _{0 . 4} The specimen of S ₁₃ obtained Seebeck coefficient values of 136-169 μVK ^{- 1} and Cu ₁₂ Sb ₃ at 723K _{. 7} Si _{0 . 3} S ₁₃ showed the highest Seebeck coefficient value of 178 μVK ^-1 .

Experimental Example 2-4: Electrical conductivity analysis of Example 2

14 is Cu ₁₂ Sb _{4 - y} Si _y S ₁₃ of The temperature dependence of electrical conductivity (σ) was shown. In the specimen with Si content y ≥ 0.2, the electrical conductivity increased with temperature, then showed a peak value at 623 K and then decreased slightly, indicating a transition characteristic from a non-degenerate semiconductor to a degenerate semiconductor (metallic). However, in the specimen with y ≥ 0.3, non-degenerate semiconductor behavior was observed in which the electrical conductivity increased with temperature. As the Si content increased at a constant temperature, the electrical conductivity decreased. Cu ₁₂ Sb _{3 . 9} Si _{0 .} The specimen of ₁ S ₁₃ obtained the highest electrical conductivity value as (2.8-3.5)Х10 ⁴ Sm ^-1 in the temperature range of 323-723K.

Experimental Example 2-5: Analysis of output factors of Example 2

Cu ₁₂ Sb _{4 - y} Si _y S ₁₃ The output factors of are shown in FIG. 15 . The output factor is calculated as PF = α ² σ and is proportional to the Seebeck coefficient and electrical conductivity. However, since the Seebeck coefficient is inversely proportional to the carrier concentration and the electrical conductivity is proportional to the carrier concentration, the output factor has a complementary relationship to these two parameters. As the temperature increased, the output factor increased. This is a result of the temperature dependence of the Seebeck coefficient and electrical conductivity. When the Si content was increased at a certain temperature, the output factors tended to decrease (especially at high temperatures). This means that the decrease in electrical conductivity was more dominant in the decrease in the output factor than the increase in Seebeck coefficient by Si doping. Cu ₁₂ Sb _{3 . 9} Si _{0 .} In the case of ₁ S ₁₃ specimen, the Seebeck coefficient had the lowest value, but due to the influence of the highest electrical conductivity value, the maximum output factor value was 0.86 ^{mWm -1} K ^-2 at 723 K.

Experimental Example 2-6: Thermal conductivity analysis of Example 2

16 is Cu ₁₂ Sb _{4 - y} Si _y S ₁₃ It is a graph showing the thermal conductivity (κ) of . Thermal conductivity is expressed as the sum of electronic thermal conductivity (κ _E ) and lattice thermal conductivity (κ _L ), and electronic thermal conductivity was calculated according to Wiedemann-Franz law (κ _E = LσT, L: Lorenz number), and Lorenz number is It was calculated using the formula L[10 ^-8 V ² K ^-2 ] = 1.5 + exp(-│σ│/116). As shown in (b) of FIG. 16 , the electronic thermal conductivity decreased as the Si content increased, and showed a tendency to increase as the temperature increased, but the lattice thermal conductivity was 0.54-at 323-723 K without a specific tendency of the Si content and temperature dependence. 0.72 Wm ^-1 K ^-1 was shown. As shown in (a) of FIG. 16 , the total thermal conductivity increased with an increase in temperature in the specimen with y ≤ 0.3 at 623 K, and in the specimen with y = 0.4. Therefore, except at 723 K, when the Si content increased at a certain temperature, the overall thermal conductivity decreased. Cu ₁₂ Sb _{3 . 6} Si _{0 .} The specimen of ₄ S ₁₃ showed a minimum thermal conductivity of ^0.74-0.92 Wm ^-1 K ^-1 in the temperature range of 323-623 K, and Cu ₁₂ Sb ₃ at 723 _{K. 7} Si _{0 . 3} S ₁₃ showed the lowest thermal conductivity of 0.85 Wm ^-1 K ^-1 .

17 is Cu ₁₂ Sb _{4 - y} Si _y S ₁₃ The temperature dependence of the Lorenz number of In all specimens, the Lorenz number decreased as the temperature increased, and the Lorenz number decreased when the Si content increased at a constant temperature. Lorenz number theoretically has a range of (1.45-2.44) × 10 ^-8 V ² K ^-2 , and the smaller the value, the more non-degenerate semiconductor behavior. As shown in FIG. 17 and Table 2, the Lorenz number of Si-doped Tetrahedrite was (1.81-1.86) × 10 ^-8 V ² K ^-2 at 323 K.

Experimental Example 2-7: Thermoelectric figure of merit (ZT) analysis of Example 2

18 shows the dimensionless figure of merit (ZT ₌ α ² σTκ ^{- 1} ) of Cu ₁₂ Sb _{4 -y} Si _y S ₁₃ . Due to the increase of the output factor and the maintenance of low thermal conductivity, the ZT value increased as the temperature increased. The change in ZT according to the Si doping content was not large, but when y ≥ 0.3, ZT decreased at high temperature. Cu ₁₂ Sb ₃ showing relatively low thermal conductivity (0.90 Wm ^-1 K ^-1 ) and high power factor (0.79 mWm ^-1 K ^-2 ) _{. 8} Si _{0 . 2} For the _S13 specimen, the maximum dimensionless figure of merit of ZT = 0.63 was obtained at 723 K. Through this embodiment, Si-doped tetrahydrite that does not use toxic and rare elements such as As, Te, and Bi can be synthesized and molded in a relatively short time through a method using mechanical alloying and hot compression molding. It was confirmed that there is a useful process.

Example 3: Sn-doped tetrahedrite Cu ₁₂ Sb _y Sn _4-y S ₁₃

Sn-doped tetrahedrite Cu ₁₂ Sb _y Sn _{4 - y} S ₁₃ (y = 0.1, 0.2, 0.3 and 0.4) were synthesized using mechanical alloying. Raw material powder Cu (purity 99.9%, < 45㎛, Kojundo), Sb (purity 99.999%, < 150㎛, Kojundo), Sn (Purity 99.999%, < 35㎛, Kojundo), S (purity 99.99%, < 75㎛) , Kojundo) was weighed according to the chemical composition, and the weight ratio (BPR) of the powder to the ball was set to 20, and a 5 mm diameter STS ball and STS jar were used in an Ar atmosphere using a planetary mill (Fritsch Pulverisette5) at a rotation speed of 350 rpm. was subjected to mechanical alloying for 24 hours. The synthesized powder was charged into a graphite mold (inner diameter 10 mm) and vacuum hot compression molding (HP) was performed at 723 K for 2 h at a pressure of 70 MPa.

[Experimental Example 3]

Experimental Example 3-1: X-ray diffraction pattern analysis of Example 3

The results of X-ray diffraction analysis of Cu ₁₂ Sb _{4 -y} Sn _y S ₁₃ ( _y = 0.1, 0.2, 0.3, 0.4) produced by mechanical alloying and hot compression molding are shown in FIG. 19 . No phase change according to the doping content was observed and it appeared as a tetrahedrite single phase without secondary phase (ICDD PDF#024-1318).

Meanwhile, the lattice constants calculated through Rietveld refinement are shown in Table 3 below.

Composition Relative
density
[%] Lattice
constant
[nm] Lorenz number
[10 ^-8 V ² K ^-2 ] Nominal Actual Cu ₁₂ Sb _3.9 Sn _0.1 S ₁₃ Cu _13.59 Sb _3.07 Sn _0.10 S _12.39 99.9 1.0348 1.82 Cu ₁₂ Sb _3.8 Sn _0.2 S ₁₃ Cu _12.97 Sb _2.81 Sn _0.20 S _13.01 100.0 1.0359 1.79 Cu ₁₂ Sb _3.7 Sn _0.3 S ₁₃ Cu _13.48 Sb _2.68 Sn _0.31 S _12.53 99.5 1.0357 1.73 Cu ₁₂ Sb _3.6 Sn _0.4 S ₁₃ Cu _13.00 Sb _2.57 Sn _0.39 S _13.04 99.7 1.0364 1.59

The lattice constant of Cu ₁₂ Sb _{4 -y} Sn _y S ₁₃ shown in Table 3 increased from _1.0348 to 1.0364 nm as the Sn content increased. According to previous studies, it has been reported that Sn ⁴⁺ is doped into Cu in the absence of other divalent transition metals in the Tetrahedrite system ^. Therefore, Sn ²⁺ is highly likely to be substituted at ^the Sb ³⁺ site, ^and as Sn ⁴⁺ is substituted ^with Cu ^{+ and} Cu ²⁺ , the lattice constant seems to increase as the Sn content increases.

Experimental Example 3-2: Analysis of the shape and composition of Example 3

20 is a SEM photograph of a fracture _surface of Cu ₁₂ Sb _{4 -y} Sn _y S ₁₃ . The difference in microstructure of the fracture surface according to the Sn content was not observed, and pores were hardly observed. Referring to FIG. 21 , the EDS element mapping indicates that all elements are homogeneously well distributed. As shown in Table 3 above, all specimens obtained a high relative density of 99.5-100.0%, and there was no significant difference between the actual composition and the nominal composition. Therefore, in this embodiment, it can be seen that the synthesis through mechanical alloying and uniform and dense sintering through hot compression molding were successfully achieved.

Experimental Example 3-3: Electrical conductivity analysis of Example 3

22 is _related to the electrical conductivity of Cu ₁₂ Sb _{4 -y} Sn _y S ₁₃ . Semiconductor electrical conductivity is affected by carrier concentration and mobility. Electrical conductivity is expressed as σ = neμ (n: carrier concentration, e: electronic charge, μ: mobility,). In the case of the y = 0.1 specimen, as the temperature increased, it showed almost no temperature dependence up to 623 K, but showed a slight decrease at a temperature of 623 K or higher, indicating degenerate semiconductor behavior. In the case of the y = 0.2-0.3 specimen, the electrical conductivity slightly increased as the temperature increased, and the y = 0.4 specimen showed a strong positive temperature dependence, showing non-degenerate semiconductor behavior. Therefore, it is judged that the electrical conduction mechanism is transferred from a degenerate semiconductor to a non-degenerate semiconductor by Sn doping in the tetrahedrite. As the content of Sn increased at a certain temperature, the electrical conductivity decreased. Therefore, at a temperature of 323-723 K, Cu ₁₂ Sb _{3 . 9} Sn _{0 .} The electrical conductivity of ₁ S ₁₃ was (2.24-2.40)×10 ⁴ Sm ^-1 , indicating the highest electrical conductivity.

Experimental Example 3-4: Analysis of Seebeck coefficient of Example 3

= (8/3)π²k_B ²m^*Te^-1h^-2(π/3n)^2/3 (k_B: Boltzmann constant, m^*: effective carrier mass, h: Planck constant,)로 나타내어진다. 온도가 증가함에 따라 y ≤ 0.3인 경우 Seebeck coefficient는 증가하였지만 y = 0.4인 경우 온도 증가에 따라 감소하는 경향을 나타내었다. P-type 반도체는 특정온도(진성 전이 온도) 이상이 되면 캐리어 농도가 급격히 증가하게 되면서 제백 계수가 감소 한다. 이것은 도 22에서 전기전도도의 온도 의존성과 관련이 있다. 따라서 Cu₁₂Sb_3.6Sn_0.4S₁₃ 의 시편은 323 K 이하에서 진성 전이가 발생한 것으로 볼 수 있다. 일정온도에서 Sn 함량이 증가함에 따라 제벡계수는 증가하였다. Cu₁₂Sb_{3 . 6}Sn_{0 . 4}S₁₃는 323-723 K에서 270-238 μVK^-1으로 최대 제벡계수를 나타냈다. 23 _shows the Seebeck coefficient of Cu ₁₂ Sb _{4 -y} Sn _y S ₁₃ . The Seebeck coefficients of all the specimens showed positive values, which means that the main carriers were p-type semiconductors with holes. Seebeck coefficient in p-type semiconductor is

= (8/3)π ² k _B ² m ^* Te ^-1 h ^-2 (π/3n) ^2/3 (k _B : Boltzmann constant, m ^* : effective carrier mass, h: Planck constant,) . As the temperature increases, the Seebeck coefficient increases when y ≤ 0.3, but decreases with the increase in temperature when y = 0.4. In P-type semiconductors, when a certain temperature (intrinsic transition temperature) is exceeded, the carrier concentration rapidly increases and the Seebeck coefficient decreases. This is related to the temperature dependence of the electrical conductivity in FIG. 22 . Therefore, in the specimen of Cu ₁₂ Sb _3.6 Sn _0.4 S ₁₃ , it can be considered that the intrinsic transition occurred at 323 K or less. As the Sn content increased at a constant temperature, the Seebeck coefficient increased. Cu ₁₂ Sb _{3 . 6} Sn _{0 . 4} S ₁₃ showed the maximum Seebeck coefficient of 270-238 μVK ^-1 at 323-723 K.

Experimental Example 3-5: Analysis of the output factor of Example 3

식 으로 계산된다. 따라서 도 22의 전기 전도도와 도 23의 제백계수의 온도의존성 결과로부터 출력인자는 온도가 증가함에 따라 증가하였다. 제백계수와 전기전도도는 모두 캐리어 농도의 영향을 받기 때문에 Sn의 함량이 증가함에 따라 캐리어 농도(전기전도도)가 감소하여 출력인자는 감소하는 결과를 보여주었다. 본 실험예에서의 최대 출력인자는 Cu₁₂Sb_{3 . 9}Sn_{0 . 1}S₁₃가 323-723 K의 온도에서 0.38-0.73 mWm^-1K^{- 2}으로 가장 높았다. 24 _shows an output factor of Cu ₁₂ Sb _{4 -y} Sn _y S ₁₃ . The output argument is PF =

is calculated as Therefore, from the temperature dependence of the electrical conductivity of FIG. 22 and the Seebeck coefficient of FIG. 23, the output factor increased as the temperature increased. Since both the Seebeck coefficient and the electrical conductivity are affected by the carrier concentration, as the Sn content increased, the carrier concentration (electrical conductivity) decreased and the output factor decreased. The maximum output factor in this experimental example is Cu ₁₂ Sb _{3 . 9} Sn _{0 . 1} S ₁₃ was the highest at a temperature of 323-723 K ^with 0.38-0.73 mWm ^-1 K ^-2 .

Experimental Example 3-6: Thermal conductivity analysis of Example 3

식을 사용하여 계산하였다.25 _shows the thermal conductivity of Cu ₁₂ Sb _{4 -y} Sn _y S ₁₃ . The lattice thermal conductivity (κ _L ) was obtained by subtracting the electronic thermal conductivity (κ _E ) from the total thermal conductivity (κ) from the κ = κ _E + κ _L relationship, and the electronic thermal conductivity was determined by the Wiedemann-Franz law (κ _E = LσT, L: Lorenz number) was used. At this time, Lorenz number [10 ^-8 V ² K ^{- 2} ] is

It was calculated using the formula.

Figure 25 ₍ a) is Cu ₁₂ Sb _{4 -y} Sn _y S ₁₃ showed the overall thermal conductivity of As the temperature increased, the thermal conductivity increased, but decreased above 623 K under the influence of the lattice thermal conductivity. At a constant temperature, the thermal conductivity decreased with increasing Sn content. Cu ₁₂ Sb _{3 . 6} Sn _{0 . 4} S ₁₃ showed the lowest thermal conductivity of ^0.49-0.60 Wm ^-1 K ^-1 at 323-723 K. It was confirmed that the thermal conductivity decreased as Sn was substituted at the Sb site. This scatters high-speed acoustic phonons in the CuS ₃ triangular plane, a tetrahedrite substructure, and mixes with acoustic dispersion, reducing thermal conductivity.

As shown in (b) of FIG. 25 , the electronic thermal conductivity increased as the temperature increased. As the Sn content increased at 323-723 K, the electronic thermal conductivity decreased, Cu ₁₂ Sb _{3 . 6} Sn _{0 . 4} S ₁₃ showed the minimum electronic thermal conductivity as 0.002-0.057 Wm ^-1 K ^-1 .

26 is Cu ₁₂ Sb _{4 - y} Sn _y S ₁₃ It is the result showing the Lorentz constant of . In general, the Lorentz constant has a value of (1.45-2.44)×10 ^-8 V ² K ^{- 2} , and a larger Lorentz constant means a degenerate semiconductor behavior and a smaller Lorentz constant means a non-degenerate semiconductor behavior. In this experimental example, at y ≤ 0.3, the Lorentz constant decreases as the temperature increases, and it is determined that the conduction characteristics of the degenerate semiconductor are transferred to the non-degenerate semiconductor. Cu ₁₂ Sb _{3 . 9} Sn _{0 . 1} S ₁₃ showed the highest Lorentz constant value as (1.82-1.71)×10 ^-8 V ² K ^-2 at 323-723 K, and Cu ₁₂ Sb _{3 . 6} Sn _{0 . 4} S ₁₃ showed the lowest Lorentz constant from 323-723 K to (1.59-1.62) × 10 ^-8 V ² K ^-2 as the Lorentz constant increased as the temperature increased.

Experimental Example 3-7: Analysis of the thermoelectric figure of merit (ZT) of Example 3

27 shows the dimensionless figure of merit ₍ ZT) of Cu ₁₂ Sb _{4 -y} Sn _y S ₁₃ . As the temperature increased, the ZT increased due to the increase of the output factor. Although the thermal conductivity decreased when the Sn content was increased, the decrease in the output factor dominated and the ZT decreased. In ZT, Cu ₁₂ Sb _3.9 Sn _0.1 S ₁₃ showed the maximum value of 0.66 at 723 K.

Features, structures, effects, etc. described in the above-described embodiments are included in at least one embodiment of the present invention, and are not necessarily limited to one embodiment. Furthermore, the features, structures, effects, etc. illustrated in each embodiment can be combined or modified for other embodiments by those of ordinary skill in the art to which the embodiments belong. Accordingly, the contents related to such combinations and modifications should be interpreted as being included in the scope of the present invention.

In addition, although the embodiments have been described above, these are merely examples and do not limit the present invention, and those of ordinary skill in the art to which the present invention pertains are exemplified above in a range that does not depart from the essential characteristics of the present embodiment. It can be seen that various modifications and applications that have not been made are possible. For example, each component specifically shown in the embodiments may be implemented by modification. And differences related to these modifications and applications should be construed as being included in the scope of the present invention defined in the appended claims.

Claims (7)

A thermoelectric material in which Cu ₁₂ Sb ₄ S ₁₃ is doped with at least one selected from the group consisting of Ge, Si and Sn. According to claim 1,
The thermoelectric material is Cu ₁₂ Sb _4-y M _y S ₁₃ ,
Here, 0 < y ≤ 1.5, and M is at least one selected from the group consisting of Ge, Si, and Sn thermoelectric material. Preparing a raw material by weighing it according to the compositional formula Cu ₁₂ Sb _{4 - y} M _y S ₁₃ (where 0 < y ≤ 1.5, and M is at least one selected from the group consisting of Ge, Si and Sn);
synthesizing powder by mechanically alloying the prepared raw material; and
A method of manufacturing a thermoelectric material comprising the step of hot pressing the powder (Hot Pressing, HP). 4. The method of claim 3,
In the step of synthesizing the powder, the method for producing a thermoelectric material, characterized in that the prepared raw material is performed under the conditions of 200 rpm to 2000 rpm. 4. The method of claim 3,
In the step of synthesizing the powder, the method for producing a thermoelectric material, characterized in that ball milling the prepared raw material for 1 hour to 100 hours. 4. The method of claim 3,
The step of hot compression molding is a method of manufacturing a thermoelectric material, characterized in that it is performed in a temperature range of 473 K to 873 K. 4. The method of claim 3,
The method of manufacturing a thermoelectric material, characterized in that the hot compression molding is performed in a pressure range of normal pressure to 100 MPa.

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