Recent advances in thermoelectric materials

Progress in Materials Science 83 (2016) 330–382

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Progress in Materials Science journal homepage: www.elsevier.com/locate/pmatsci

Recent advances in thermoelectric materials Chhatrasal Gayner a, Kamal K. Kar a,b,⇑ a b

Advanced Nanoengineering Materials Laboratory, Materials Science Programme, Indian Institute of Technology Kanpur, Kanpur 208016, India Advanced Nanoengineering Materials Laboratory, Department of Mechanical Engineering, Indian Institute of Technology Kanpur, Kanpur 208016, India

a r t i c l e

i n f o

Article history: Available online 7 July 2016

a b s t r a c t Thermoelectric materials are crucial in renewable energy conversion technologies to solve the global energy crisis. They have been proven to be suitable for high-end technological applications such as missiles and spacecraft. The thermoelectric performance of devices depends primarily on the type of materials used and their properties such as their Seebeck coefficient, electrical conductivity, thermal conductivity, and thermal stability. Classic inorganic materials have become important due to their enhanced thermoelectric responses compared with organic materials. In this review, we focus on the physical and chemical properties of various thermoelectric materials. Newly emerging materials such as carbon nanomaterials, electronically conducting polymers, and their nanocomposites are also briefly discussed. Strategies for improving the thermoelectric performance of materials are proposed, along with an insight into semiconductor physics. Approaches such as nanostructuring, nanocomposites, and doping are found to enhance thermoelectric responses by simultaneously tuning various properties within a material. A recent trend in thermoelectric research shows that high-performance thermoelectric materials such as inorganic materials and carbon nanomaterials/electronically conducting polymer nanocomposites may be suitable for power generation and energy sustainability in the near future. Ó 2016 Elsevier Ltd. All rights reserved.

Contents 1. 2. 3.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Historical background of thermoelectrics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Terminologies of thermoelectrics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1. Thermoelectric figure of merit . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2. Power conversion efficiency . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

331 333 334 334 335

Abbreviations: a, Seebeck coefficient; c, acoustic phonon Grüneisen parameter; hD, Debye temperature; j, total thermal conductivity; jlattice, lattice thermal conductivity; jC, je, charge-carrier thermal conductivity; l, mobility; g, efficiency; q, density; r, electrical conductivity; a, lattice parameter; Cp, specific heat; h, Planck’s constant; kB, Boltzmann constant; Tcold, cold-end temperature; Thot, hot-end temperature; ZT, figure of merit; ASSET, atomic-scale structural engineering of thermoelectrics; CVD, chemical vapor deposition; CNTs, carbon nanotubes; DOS, density of states; DMSO, dimethyl sulfoxide; EG, ethylene glycol; fH, full-Heusler; HP, hot pressing; HH, half-Heusler; LAST, AgPbmSbTe2+m; LASTT, AgPbmSnnSbTe2+m+n; MWCNTs, multiwalled carbon nanotubes; NPs, nanoparticles; PEDOT, poly(3,4-ethylenedioxythiophene); PEDOT:Tos, PEDOT:tosylate; PGEC, phonon glass electron crystal; PTH, polythiophene; PSS, poly(styrenesulfonate); PANI, polyaniline; PVDF, polyvinylidene difluoride; QDSL, quantum dot superlattice; SEM, scanning electron microscope; SWCNTs, single-walled carbon nanotubes; SPS, spark plasma sintering; Tos, tosylate; vol%, volume percent. ⇑ Corresponding author at: Advanced Nanoengineering Materials Laboratory, Materials Science Programme, Indian Institute of Technology Kanpur, Kanpur 208016, India. E-mail address: [email protected] (K.K. Kar). http://dx.doi.org/10.1016/j.pmatsci.2016.07.002 0079-6425/Ó 2016 Elsevier Ltd. All rights reserved.

C. Gayner, K.K. Kar / Progress in Materials Science 83 (2016) 330–382

4.

5.

6.

7.

3.3. Carrier concentration . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.4. Seebeck coefficient . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.5. Effective mass . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.6. Electrical conductivity. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.7. Thermal conductivity . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Complex thermoelectric materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.1. Classification of thermoelectric materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2. Syntheses of thermoelectric materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.1. Melt and growth. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.2. Levitation or arc melting . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.3. Mechanical alloying . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.4. Single-crystal growth . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.5. Hydrothermal or solvothermal . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.6. Low-temperature aqueous chemical route . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.7. Polyol process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.8. Sol–gel process . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.9. Microwave synthesis . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2.10. Other methods . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Strategies for enhancing the performance of thermoelectrics. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.1. Doping . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2. Alloying . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3. Nanostructuring. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.4. Nanocomposites and nanoinclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.5. Thin films and superlattices . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Performance evaluation of thermoelectric materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.1. Effect of material properties on performance . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.1.1. Crystal structure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.1.2. Microstructure/Densification . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.1.3. Grain orientation and size . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2. Thermoelectric properties of inorganic materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.1. Skutterudites. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.2. Telluride-based materials: PbTe . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.3. Bi2Te3. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.4. AgSbTe2 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.5. AgPbmSbTe2+m (LAST) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.6. AgPbmSnnSbTe2+m+n (LASTT). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.7. Rare earth chalcogenides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.8. Replacement of Tellurium: PbSe/Bi2S3 . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.9. Copper ion liquid-like thermoelectric Cu2Se. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.10. Si–Ge alloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.11. HH alloys. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.12. Clathrates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.2.13. Other emerging systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.3. Thermoelectric properties of electronically conducting polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.3.1. Poly(3,4-ethylenedioxythiophene). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.3.2. PEDOT:PSS. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.3.3. Poly(2,7-carbazole) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.3.4. Polyaniline . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.4. Thermoelectric properties of carbon nanomaterials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.4.1. Carbon nanotubes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.4.2. Graphene. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.5. Thermoelectric properties of carbon nanomaterial/polymer nanocomposites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.5.1. CNT/electronically conducting polymer nanocomposite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6.5.2. Graphene/electronically conducting polymer composite . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Perspective and concluding remarks. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Acknowledgment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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335 336 337 338 338 340 340 340 340 341 341 341 341 341 341 341 342 342 342 343 346 346 347 349 349 349 350 350 350 351 351 354 356 357 357 358 359 360 360 361 362 364 365 367 367 368 368 369 369 370 370 372 375 376 376 376 376

1. Introduction Humankind is facing a serious increasing demand for energy. With the decrease in fossil fuels, renewable energy sources will be essential for resolving the power crisis in the future. In this respect, energy conversion technologies such as solar cells and fuel cells are highly relevant, although their global commercialization is limited by their poor efficiency, high cost, and

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Fig. 1. Schematic represents the global energy conversion (in%) by various technologies [5,10,11].

8%

4%

12%

17%

18%

PbTe Bi Te3

2% 2%

2

SiGe Skutterudite 19% Clathrate Half heusler alloy Other emerging inorganic materials Conducting polymers CNT/graphene- polymer composite

18%

Fig. 2. Schematic shows the various materials used in thermoelectric research and development and their individual contribution [8,9].

poor long-term stability. Thermoelectric devices are used for the direct conversion of heat to electricity. In principle, these devices can use any thermal source including solar and waste heat. Therefore, thermoelectric materials play a key role in the development of sustainable energy-efficient technologies. Furthermore, these thermoelectric devices do not have any moving parts and are thus free of vibration and noise [1,2]. Current thermoelectric materials exhibit a conversion efficiency of 5–20%, which can be further enhanced with appropriate strategies such as doping, alloying, and nanostructuring. In the case of thermoelectric materials, the thermoelectric figure of merit (ZT) can be defined as follows:

ZT ¼

a2 r j

ð1Þ

where r is the electrical conductivity, a is the Seebeck coefficient, and j is the thermal conductivity of the material. For ideal thermoelectric materials, ZT should be P1 to obtain a conversion efficiency of >10% [1,3–7]. In the past 15 years, the significant growth in the thermoelectric energy conversion sector is reflected in the increase in related annual publications from 500 to 2500 [8,9]. The contribution of thermoelectric technologies along with other established energy conversion technologies is shown in Fig. 1. The fundamental principles of thermoelectricity have been established over the past 200 years. The development of superior thermoelectric materials requires an in-depth understanding of the electron and phonon transport phenomena in such materials. Inorganic materials are classic materials used in thermoelectric research. Some examples include skutterudites, telluride-based materials (e.g., PbTe, Bi2Te3, etc.), rare earth chalcogenides (e.g., La3xTe4), copper ion liquid-like materials (e.g., Cu2Se), Si–Ge alloys, half-Heusler (HH) alloys, and clathrates. Recently, new classes of materials such as carbon nanomaterials, electronically conducting polymers, and carbon nanomaterial/polymer nanocomposites have been proposed. The percentage contribution of these materials to thermoelectric technologies is presented in Fig. 2. An ideal thermoelectric material must exhibit (i) high electrical conductivity, (ii) high Seebeck coefficient, and (iii) low thermal conductivity. The high Seebeck coefficient ensures a large potential/thermovoltage, the high electrical conductivity is needed to minimize the Joule heating effect, and the low thermal conductivity is needed to create a large temperature gradient [1,3,6]. Proper control of these parameters for a single material is a challenging task. By lowering the lattice thermal conductivity and by optimizing the electronic thermal conductivity, low thermal conductivity can be obtained [1–4].

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Parameters such as electrical conductivity, Seebeck coefficient, and thermal conductivity within a single material can be controlled with various strategies such as phonon glass electron crystal (PGEC) [12], doping [13–16], energy filtering, resonant states, multiple-band conduction mechanism and convergence of electronic bands [13–17], scattering from nanoscale endotaxial precipitation and mesoscale grain boundaries, atomic-scale alloy scattering [18], nanostructuring [19–25], quantum confinement, superlattices [26–41], and nanocomposites [42]. The dispersion of nanophase inclusions (secondary phases) is an effective method of reducing the lattice thermal conductivity. Nanostructuring can significantly reduce the thermal conductivity without causing a significant increase in electrical conductivity. Either one or a combination of these strategies can be used to achieve the required parameter values. Several studies have been conducted to develop high-performance thermoelectric devices by incorporating nanostructures, improving processing techniques, and eliminating the problems of electrical and thermal transport [1,4,43,44]. Apart from thermoelectric materials, contact materials also play a key role in thermoelectric devices. The semiconductor–metal junction exhibits a greater electrical resistance than a metal–metal junction, which affects the performance of the device. Sterzel demonstrated a thermally stable semiconducting alloy for use in soldering [45]. The contact resistance can be reduced with the use of a barrier layer consisting of borides, nitrides, carbides, phosphides, or silicides. These materials can be used as contact materials, to reduce the resistance at the semiconductor–metal junction [45]. In this review, we discuss the synthesis, properties, and characterization of various thermoelectric materials including inorganic materials, organic materials, and nanocomposites. The physics underlying the thermoelectric responses such as material transport and semiconductor physics is discussed using the available literature and patents. The article focuses on the dependence of thermoelectric properties of various materials on the synthesis process, crystal structure, grain size and grain orientation, morphology, etc. Furthermore, the potential methods of enhancing thermoelectric properties are discussed. This review provides insights into the past, present, and future of thermoelectric materials.

2. Historical background of thermoelectrics 1821–1909: In 1821, the Seebeck effect was first observed. Later, in 1834, the Peltier effect was discovered. In 1851, Gustav Magnus stated that the potential difference produced between the two junctions of a thermocouple is directly proportional to the applied temperature difference. In 1856, the Thomson effect was proved experimentally [1,4]. 1910–1949: Altenkirch first introduced the term ‘‘thermoelectric figure of merit”. He suggested that a thermoelectric material exhibit a large Seebeck coefficient, a high electrical conductivity to minimize Joule heating, and a low thermal conductivity to maintain a large thermal gradient. Bloch showed that the thermal vibration of the lattice and the addition of impurities affect the electrical conductivity of metals. Wilson suggested that, for semiconductors and insulators, the conductivity should be dependent on temperature. At lower temperatures, the conductivity decreases because of low carrier concentration [1,3,46]. 1950–1959: Ioffe stated that the thermal conductivity of semiconductors is a function of atomic weight. Generally, elements with large atomic weights have low thermal conductivity (although low thermal conductivity is also reported in low-density materials) [1,4]. Horovitz proposed a model demonstrating low-temperature deviations from the lattice equilibrium [1,4,46]. Goldsmid studied the variation in electrical conductivity with the crystal structure and electron mobility [4]; the mobility was found to affect the thermoelectric properties of a material, and the thermal conductivity was dependent on the atomic weight. Goldsmid proposed that the ratio of electron mobility to thermal conductivity is a function of atomic weight. Ioffe hypothesized that alloying semiconductors with isomorphous materials can reduce the thermal conductivity without affecting its electrical conductivity. Thus, modified semiconductor alloys were used for thermoelectric applications. In this context, Ioffe proposed ZT as a measure to quantify the thermoelectric performance of materials [1,4,46]. 1960–1969: Chasmar and Stratton introduced a material parameter (fit parameter) as a function of carrier mobility, effective mass, doping, temperature, and thermal conductivity [3]. Using Ioffe’s model, the effect of bandgap on the ZT can be demonstrated. A large bandgap leads to high thermal conductivity and low carrier mobility. Littman and Davidson proposed the lack of an upper limit for ZT based on irreversible thermodynamics. 1970–2015: Dresselhaus and Hicks stated that layering highly anisotropic thermoelectric materials (such as superlattices) should increase ZT, although the presence of different orientations/interfaces reduces the thermal conductivity because of phonon scattering. Their argument was further modified by Sofo and Mahan, such that the alternating barrier layers of superlattices have finite thermal conductivity and tunneling probability in the quantum wells. They also suggested that the quantum mixing in the well changes the density of states (DOS) from 2D to 3D and predicted an increase in ZT, but below that predicted by Dresselhaus and Hicks [26,27,47]. Bulusu and Walker proposed a nonequilibrium Green’s function method to determine the thermoelectric performance of silicon nanofilms, nanowires, and Si/Ge/Si superlattices. The dispersion relation for each sub-band is used to model the quantum effects. The effects of confinement and scattering on the electrical conductivity and Seebeck coefficient in superlattices were studied [46]. Among bulk thermoelectric materials, in addition to commercially available materials (e.g., PbTe, Bi2Te3, etc.), new classes of materials have been proposed with the aim of reducing the thermal conductivity and increasing the power factor (a2r) [12,19,23–25]. Slack proposed a model with atoms that rattle inside a cage-like structure and introduced the novel concept

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of ‘‘PGEC” [12]. Heremans, Snyder, and others focused on doping in semiconductors and its effectiveness in improving the thermoelectric response [13–15]. From bandgap engineering, new concepts have emerged, such as energy filtering, pinning of energy levels, multiple-band conduction mechanisms, and convergence of electronic bands [13–17]. Nanocomposites were first introduced as novel thermoelectric materials. Kanatzidis et al. proposed a panoscopic approach combining band engineering and nanostructuring to produce high-performance bulk thermoelectrics [18]. Tritt, Poon, Poudeu, Ren and other research groups attained high ZT measures for both low-temperature (e.g., Bi2Te3) and hightemperature materials (e.g., HH alloys) with the same combination. The following approaches have been introduced: atomic-scale structural engineering of thermoelectrics (ASSET) [48], a panoscopic approach involving hierarchical scattering of phonons, via atomic-scale lattice disorders and nanoscale endotaxial precipitates to mesoscale grain boundaries [18]; nanostructure engineering; band structure engineering; and a synergistic approach of high-mobility electron injection, energy filtering, and boundary scattering [49]. Strategies for enhancing the thermoelectric response of the materials via nanostructuring, doping, use of nanocomposites, etc. have also been introduced. The present study focuses on developing flexible thermoelectric materials using electronically conducting polymers. For high-performance and lightweight thermoelectric devices, various carbon nanomaterials/electronically conducting polymer nanocomposites are also being developed [50].

3. Terminologies of thermoelectrics 3.1. Thermoelectric figure of merit In a thermoelectric device, the charge carriers are transported by the formation of p–n junctions by p- and n-type materials, with electrons/holes acting as a working ‘‘fluid” (Fig. 3). The applied temperature gradient generates gradients of charge carriers, which diffuse from the hot side to the cold side, in turn producing an electrostatic potential (DV). This thermoelectric potential is directly proportional to the temperature difference, DT ðDV ¼ a  DTÞ, where a is the Seebeck coefficient [1–4]. The diffusion of charge carriers and the flow of charge carriers (mobility) are the parameters determining the thermoelectric transport. A dimensionless parameter, ZT, as defined by Eq. (1), is used to assess the thermoelectric transport properties of a material [1,3–5]. The aim of the present study is to enhance ZT, either by increasing the power factor or by reducing the thermal conductivity. Therefore, ideal thermoelectric materials must exhibit a large Seebeck coefficient, a high electrical conductivity, and a low thermal conductivity to produce a large potential difference across the junction, so as to minimize the Joule heating effect and to maintain a large temperature gradient, respectively [1,3,6]. Fig. 4 illustrates the ZT values of the selected materials.

Fig. 3. A thermoelectric couple operating across a differential temperature and generating electrical power. Reproduced with permission from [5].

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Fig. 4. Plot shows the ZT values of different classes of materials over a range of temperatures: PbTe–SrTe [18], AgPb18SbTe20 [19], 2% Na-doped PbTe0.85Se0.15 [15], Tix(Zr0.5Hf0.1)1xNiSn half-Heusler alloy [51,52], Al-doped PbSe [53], Cu2Se1.01 [54], Ba8Ga16Ge30 (clathrates) [55], Ba0.08Yb0.04La0.05Co4Sb12 (skutterudites) [56], melt-spun + spark plasma-sintered (SPS) Bi2Te3 [57], AgSbTe2 [58], SiGe [59], In4Se3d [60], Yb14Mn0.8Al0.2Sb11 [61], Bi0.875Ba1.25CuSeO [62], Bi-doped Mg2.16(Si0.4Sn0.6)1yBy [63], SnSe [64], Bi0.5Sb1.5Te3 [65], and organic material (inserted box) [66].

3.2. Power conversion efficiency The power conversion efficiency of a thermoelectric device can be defined as follows:

g¼

pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi DT 1 þ ZT  1 pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi T Thot 1 þ ZT þ Tcold hot

ð2Þ

The product of the Carnot efficiency (DT/Thot) and the ZT dependent quantity gives the efficiency [Eq. (2)], where Thot and Tcold denote the hot-end and cold-end temperatures, respectively. The different effects of ZT on efficiency with varying Thot are plotted in Fig. 5. Here, the conversion efficiency is 10–15% for ZT  1 and 20–25% for ZT  3 (Fig. 5) [5]. A ZT value > 2 has been achieved already with materials currently in use [18,19,30,31]; in the near future, a ZT value  3 can also be expected. For ideal thermoelectric materials, ZT should be equal to or >1 to obtain a conversion efficiency of >10% [1,3–7]. 3.3. Carrier concentration The number of charge carriers determines the electronic conductivity as well as the Seebeck coefficient. Thus, thermoelectric research requires the majority charge-carrier concentration to be optimized, to prevent bipolar conduction. Hall

Fig. 5. Plot shows the thermoelectric conversion efficiency as a function of differential operating temperature and ZT. Reproduced with permission from [5].

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Fig. 6. Plot shows the effect of carrier concentration on the ZT values of MgxPb1xTe: Na [14], PbTe1xSex [15], Pb1xMnxTe: Na [16], and La3xTe4 [69] at different temperatures. Reproduced with permission from [14–16,69].

effect measurements are used to determine the carrier concentration, mobility, and the type of charge carriers in semiconductors of both p-type and n-type [67]. The dopant concentration and the nature of the dopant further affects the chargecarrier type and carrier concentration [14,17,68]. For thermoelectric materials, the carrier concentration determines both the Seebeck coefficient and the electrical conductivity. Fig. 6 shows the effect of carrier concentration on the ZT values, which are determined by the doping elements, temperature, and charge-carrier concentration [14–16,69]. Experimentally, the carrier concentration has been found to increase with temperature, which can be tuned further by doping or alloying [14,70]. At higher temperatures, the increase in charge-carrier concentration in turn increases the electrical conductivity, with a high ZT value. The charge-carrier mobility plays a significant role in determining electrical conductivity. The mobility can be affected by doping, alloying, grain boundaries, impurity, density, and addition of dispersed secondary phases [14,15,71–73], as discussed in subsequent sections. 3.4. Seebeck coefficient The Mott relation for the Seebeck coefficient (a) is presented in Eq. (3) [13]:

a¼

p2 jB 3 q



jB T

1 dnðEÞ 1 dlðEÞ þ n dE l dE

 ð3Þ E¼EF

where T is the temperature, n(E) is the carrier density at energy E, l(E) is the mobility at energy E, EF is the Fermi energy, and q is the electronic charge. An ideal thermoelectric material must have a high Seebeck coefficient (>200 lV/K), indicating a high voltage-generating ability. Ideally, the Seebeck coefficient depends on the bandgap and the carrier concentration, and it varies as a nonlinear function of temperature. The Seebeck coefficient depends on the absolute temperature, composition, charge-carrier concentration, and crystal structure of the conductor. The variations in the Seebeck coefficient and electrical conductivity with the carrier concentration are plotted in Fig. 7. From this plot, it can be observed that the charge-carrier concentration influences both the Seebeck coefficient and the electrical conductivity, such increased carrier concentration increases the electrical conductivity and decreases the Seebeck coefficient. Hence, an optimal carrier concentration can tune the power factor; implying that degenerate semiconductors are potential candidates. The applied temperature difference generates charge carriers, that is, electrons or holes, which then diffuse from the hot side to the cold side in a thermocouple [1,4]. The criteria for measuring the Seebeck coefficient of a thermoelectric material are as follows: (i) both temperature and voltage should be measured simultaneously when the system is in a steady state, (ii) the voltage response to the temperature gradient should be linear, and (iii) temperature and voltage should be measured at the same point [1,4,74,75]. The typical values of the Seebeck coefficient for metals, semiconductors, and insulators are 10, 200, and >200 lV/K, respectively. The efficiency of a thermoelectric material depends on the Seebeck coefficient or thermopower; furthermore, a high efficiency can only be obtained with a high thermoelectric potential at a given temperature gradient [1,4]. The strategies for increasing the Seebeck coefficient include (i) increasing the energy dependence of l(E) by a scattering mechanism that is strongly dependent on the charge carriers and (ii) increasing the energy dependence of n(E) by a local increase in the DOS [13]. Fig. 8 shows the effect of various dopants and carrier concentrations on the Seebeck coefficient [76]. For example, in Na-doped PbSe, the Seebeck coefficient is reduced with an increase in carrier concentration. At 27 °C, the Seebeck coefficient for the same system is 17 lV/K with a carrier concentration of 17  1019 cm3 [76,77]. Tl dopants are found to have a high Seebeck coefficient that varies less with the carrier concentration; K-doped systems have Seebeck coefficients that vary

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Fig. 7. Plot shows the variations in the electrical properties of a thermoelectric material as a function of carrier concentration. Reproduced with permission from [2].

Fig. 8. Plot shows the effect of carrier concentration on the Seebeck coefficient of doped PbSe and other materials. Reproduced with permission from [13,15,53,69,76–81,89].

significantly (Fig. 8). This can be attributed to the tuning of the resonant energy levels created by Tl atoms via the Seebeck coefficient in a Tl–PbTe system [13]. Therefore, the nature of dopants and their concentration affect the Seebeck coefficient. 3.5. Effective mass In thermoelectric materials, the effective mass (m⁄) reflects the band curvature and doping is found to affect the effective mass. The electronic properties of thermoelectric materials are evaluated with parabolic band models, with the effective mass determining the electronic structure and transport. Non-parabolic infinite-gap semiconductors are classified by their energy- and/or temperature-dependent mass. In this context, transport in thermoelectric materials is clearly understood with varying effective mass. In heavily doped semiconductors in particular, at a given carrier concentration and temperature, the thermopower increases with effective mass [82]. Hence, the Seebeck coefficient can be expressed using constant scattering time approximation (CSTA), Eq. (4) [82]: 2

a¼

4p2 kB 2

eh

m T

 2=3 4p 3n

ð4Þ

For parabolic bands, the thermopower is proportional to m⁄. Conversely, for isotropic parabolic bands, the conductivity can be expressed in terms of m⁄ and the carrier concentration, r = ne2s/m⁄, where s is the scattering time and n is the carrier concentration. Therefore, a high Seebeck coefficient can be attained by generating a large DOS effective mass, which decreases the mobility (i.e., conductivity). Hence, a lower effective mass leads to high ZT values and a greater effective mass to low ZT values; doping and temperature can be used to tune the effective mass [83].

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3.6. Electrical conductivity The dependence of the electrical conductivity (r) of a semiconductor on carrier concentration and mobility is given by

r ¼ eðle  n þ lh  pÞ

ð5Þ

where le, lh, n, and p denote the electron mobility, hole mobility, number density of electrons, and number density of holes, respectively. Due to the large variation in these parameters, the electrical conductivity of semiconductors varies widely, as depicted in Fig. 9. Lattice and impurity scattering determine the mobility of both electrons and holes. The lattice vibrations increase with an increase in temperature, which lead to a decrease in mobility. In semiconductors, impurities affect the carrier mobility, and the ionized impurities can also produce crystal defects. Impurity scattering increases with a decrease in temperature, ultimately leading to a decrease in mobility, which is the reverse of lattice scattering. The type of material, doping, impurities, and temperature affect the electrical conductivity of thermoelectric materials [84–86]. The effect of temperature on the electrical conductivities of materials synthesized via various methods such as spark plasma sintering (SPS), hot pressing (HP), levitation melting, and zone melting is shown in Fig. 10. The samples prepared by SPS exhibit greater electrical conductivities than those synthesized by HP or cold pressing. This can be attributed to the change in the density and microstructure of the samples during the SPS process. Structural parameters such as grain size, strain, and lattice constant have also been found to affect electrical conductivity [86,87]. 3.7. Thermal conductivity In semiconductors, the net thermal conductivity is a sum of two contributions: one from charge carriers and the other from phonons. When a single type of charge carrier is predominant in a material, the total thermal conductivity (j) is the sum of the lattice thermal conductivity (jlattice) and the charge-carrier thermal conductivity (jC), that is, j = jC + jlattice [1–3]. A lower thermal conductivity can be attained by maximizing the jC/jlattice ratio. This is achieved by lowering the lattice thermal conductivity [1–4,13]. The charge-carrier thermal conductivity (jC) can be estimated from the Wiedemann–Franz law (jC = LrT), where L is the Lorenz constant (2.45  108 V2 K2 for metals and 1.5  108 V2 K2 for non-degenerate semiconductors) [71]. The lattice thermal conductivity (jlattice) can be expressed as follows:

Fig. 9. Schematic shows the room-temperature electrical conductivity of various materials [84].

Fig. 10. Effect of temperature on electrical conductivity of materials synthesized by various routes [18,53,54,86–89].

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jlattice ¼ D  Cp  q

339

ð6Þ

where D is the thermal diffusivity, Cp is the specific heat, and q is the density of the material. The thermal conductivity of a material can be measured in two ways: the steady-state method and the transient method [71]. The measurement of thermal conductivity is a powerful tool for investigating lattice defects or imperfections. Heat is transported in solids via the charge carriers (electrons or holes), phonons, electromagnetic waves, spin waves, or other excitations. The charge-carrier thermal conductivity is affected by bipolar diffusion when the carrier concentrations and mobilities of the electrons are comparable to that of holes. Eq. (7) is used to calculate the charge-carrier thermal conductivity to describe the bipolar diffusion term [71]:

#  2 "  2 kB Eg le =lh 2þ 4þ e kB T ð1 þ le =lh Þ2

je ¼ r T

ð7Þ

where je, kB, r, le, lh, T, and Eg are the thermal conductivity due to electrons, Boltzmann constant, electrical conductivity, electron mobility, hole mobility, temperature, and bandgap of the material, respectively. The thermal conductivity of materials is strongly affected by phonons, which are generated from the lattice vibrations. The lattice thermal conductivity depends on the crystal structure and lattice parameters of the material. Various factors such as lattice parameter, density of the material, and anharmonic lattice vibration determine the thermal conductivity, according to Eq. (8) [71]: 3

jlattice ¼

kB a4 qh3D 3 2 h c T

ð8Þ

where kB, h, a, q, hD, and c denote the Boltzmann constant, Planck’s constant, lattice parameter, material density, Debye temperature, and the acoustic phonon Grüneisen parameter (which is a measure of the anharmonic nature of lattice vibration), respectively. In the case of alloys, the composition also has an effect on the lattice thermal conductivity due to the variation in their lattice parameters. In the case of the Tix(Zr0.5Hf0.5)1xNiSn alloy, lattice parameters of 6.1 and 6.025 Å are observed for compositions of x  0 and  0.5, with lattice thermal conductivities of 2.5 and 2.28 W/m K, respectively. Hence, techniques such as doping or alloying can be used to tune the lattice parameter and in turn reduce the thermal conductivity [51] (see Fig. 11).

Fig. 11. (a) Plot of composition, x, versus lattice parameter of PbTe1xSex:2% Na and TiX(Zr0.5Hf0.5)1XNiSn [Inset (a): Plot of composition, x, versus lattice parameter of BaxLayYbzCo4Sb12; (b) plot of composition, x, versus lattice thermal conductivity of various alloys at different temperatures.

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Fig. 12. Effect of processing temperature on thermal conductivity of Hf0.75Zr0.25NiSn1ySby (y  0.02, 0.04) [86], CoSb3 having different particle sizes (HT1, 143 nm; HT2, 218 nm; HT3, 459 nm; and HT4, 771 nm and 1923 nm) [20,21] and zone-melted nanocomposite of Bi2Te3 [87]. Reproduced with permission from [20,21,86,87].

Processing techniques such as HP and SPS also determine the thermal conductivity. As the thermal conductivity depends on the density and microstructure of the material, samples prepared by SPS exhibit greater densification than those synthesized by HP. Hence, the net thermal conductivity of samples prepared by SPS is high as they have a higher electrical conductivity than the samples prepared by HP. Materials with smaller particle size are found to exhibit low thermal conductivities due to the effect of grain size, as smaller grains cause increased scattering of phonons (Fig. 12). The processing temperature is also found to affect the thermal conductivities of materials. The variations in the thermal conductivity of various thermoelectric materials at different temperatures are plotted in Fig. 12 [86]. From this plot, it can be observed that samples such as HT1, HT2, HT3, and HT4 exhibit different thermal conductivities when annealed at different temperatures. 4. Complex thermoelectric materials 4.1. Classification of thermoelectric materials Thermoelectric materials are widely classified as intermetallics, skutterudites, clathrates, HH, oxides, rare earth chalcogenides, Zintl-phase materials, pnicogens, nitrides, and their superlattice architectures. All of these materials are degenerate semiconductors, with each having a complex band structure. Carbon nanomaterials, such as carbon nanotubes (CNTs) and graphene, and electronically conducting polymers, such as polyaniline (PANI) and poly(3,4-ethylenedioxythiophene) (PEDOT), have also been shown to be useful in thermoelectric devices. Recently, the nanocomposites of various carbon nanomaterials and electronically conducting polymers have emerged as potential thermoelectric systems to develop flexible, next-generation thermoelectric devices; their thermoelectric performances, however, are much lower than the classical inorganic materials. 4.2. Syntheses of thermoelectric materials Thermoelectric materials are synthesized via both chemical and physical routes. As the routes determine the thermoelectric properties of the materials, the selection of the preparation route is critical. With chemical routes, one can control the particle size and particle size distribution very accurately, although these routes have several hazards. The physical routes are entirely dependent on the processing temperature and pressure, which determine the microstructure and other properties of thermoelectric materials. 4.2.1. Melt and growth The melt-and-growth process is commonly used to prepare thermoelectric alloys. In this method, mixed metal powders with the desired stoichiometric ratio are poured into a quartz or silica ampoule and sealed in vacuum to prevent the buildup of moisture and oxygen contamination. The sealed ampoule is heat-treated to obtain the final product. The heat treatment usually involves melting of the mixture and solidification followed by annealing to attain the desired microstructure. The processing temperatures are selected from the phase diagram of the compounds, as the melting temperature varies with various materials and their alloy compositions. After melting, the desired microstructure is obtained by annealing the material for a predetermined duration, or by processing the material via HP or SPS. For example, PbTe can be prepared by heating the material above its melting temperature (917 °C). The processing temperature also varies for materials doped or alloyed with other elements [14–18]. To synthesize polycrystalline Pb0.98Na0.02Te1xSex, the mixture of pure elements is melted at

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1000 °C and subsequently quenched [15]. I-doped Pb1xCrxTe [17], KxPb1xTe, KxPb1xSe, K0.02Pb0.98Te1ySey [76,78], polycrystalline Cu2xSe ingots [88], etc. are prepared in a similar manner. 4.2.2. Levitation or arc melting Levitation or arc melting is an instantaneous melting process to prepare alloy systems such as HH alloys, intermetallics, and other high-temperature materials. The materials are further processed by HP or SPS for better densification and a uniform microstructure. For example, Ti, Zr, Hf, V, Ni, Pd, Sn, and Sb are used to prepare a Tix(ZrHf)0.99xV0.01Ni0.9Pd0.1Sn0.99Sb0.01 alloy by arc melting in an Ar atmosphere; then, the prepared samples are subjected to a two-step annealing process in an evacuated quartz tube [51]. Similarly, to prepare Hf1xZrxNiSn1ySby, a stoichiometric mixture of elements is subjected to levitation melting in an Ar atmosphere and further processed by SPS at 830 °C [86]. 4.2.3. Mechanical alloying In the mechanical alloying process, elemental powders are mixed in stoichiometric ratios and ground in a high-energy ball mill. The fine powders are further processed into ingots/bars/blocks via the HP or SPS processes. For example, to synthesize Tl0.02Pb0.98Te nanopowder, Pb, Te, and Tl are ball-milled (with stainless steel balls) for 10–12 h, and the resulting powders are loaded into a graphite die and processed via HP thereafter [14]. To synthesize La3xTe4, La and Te chunks are ball-milled with stainless steel balls, and the milled powder is processed via Ar or vacuum HP sintering at 1027 °C [69]. Similar procedures are also used to prepare PbSe:Alx, PbSe:Clx, and PbSe:Ix [53]. Yu et al. prepared Cu2Se via ball milling followed by sintering at 400–700 °C aided by a conventional HP method [54]. 4.2.4. Single-crystal growth The Bridgman [90,91] and Czochralski [55,92] methods are used to synthesize single-crystalline materials. Zhang et al. prepared Ba8Au5.33Ge40.67 by the Bridgman method, following which they ground the resulting single crystal and processed it with SPS thereafter [91]. To synthesize Ba8Ga16Ge30, Ba, Ga, and Ge powders with a stoichiometric ratio of 8:16:29.5 are mixed in a glassy carbon crucible, placed in a steel autoclave under an Ar atmosphere, and heated up to 1052 °C [92]. The Czochralski method is known to be highly efficient in the preparation of single-crystalline materials of high purity [55,92]. 4.2.5. Hydrothermal or solvothermal The hydrothermal or solvothermal process is used to synthesize materials with nanodimensions, such as nanoparticles (NPs), nanotubes, nanoboxes, hollow nanospheres, and nanodendrites. It is simpler and more beneficial than other methods as the particle size in the nano-regime can be effectively controlled. Acetate precursors of various elements and an autoclave are required for this method. For example, PbTe powder is synthesized using lead acetate, Te, NaOH, and NaBH4 as the precursor materials taken in an autoclave, which is filled with ethanol/glycol/acetone mixture and kept at an elevated temperature to complete the reaction [93]. Zhao et al. synthesized Bi2Te3 nanopowder via the hydrothermal method [87]. 4.2.6. Low-temperature aqueous chemical route The low-temperature aqueous chemical route is very similar to the hydrothermal method, as both use the same kind of precursors to synthesize materials. The difference lies in the reaction occurring at a relatively low temperature in this method. For example, to synthesize PbTe, 1 mmol of Te and 2 g of NaBH4 are added to NaOH/H2O, and then 1 mmol of Pb (CH3COO)23H2O is added dropwise, leading to a purple-black solution. The reaction is allowed to complete, the system is cooled to room temperature, and the resulting gray powder is obtained [93]. 4.2.7. Polyol process This method is used to prepare PbTe, Ni2Te3, and Cu7Te5 under microwave irradiation. Acetates of Pb, Ni, Cu, or other elements are used as the precursors. The selected precursor is dissolved in ethylene glycol (EG), to which an elemental powder of Te (or any other desired element) is added. A solid-state moisture-assisted reaction takes place under microwave irradiation, yielding the desired material in its nanocrystalline form [94]. 4.2.8. Sol–gel process The sol–gel process is a versatile technique for preparing metal oxides, chalcogenides, and semiconductors. To synthesize NaCo2O4, a solution of metal nitrate is prepared by dissolving NaNO3 and Co(NO3)3 in deionized water. The resulting aqueous solution is a mixture of metal nitrates with atomic ratios of Na:Co  1.15:2 and urea [95]. PbTe gels and aerogels with nanodimensions are prepared using lead acetate trihydrate (1.317 g) with oleic acid (3 mL) and 1-octadecene (6 mL). The solution is initially heated at a temperature of 170 °C for 30 min in an inert atmosphere. Then it is mixed with a solution of trioctylphosphine telluride, and subsequently the entire solution is quenched in a cold water bath. The NPs are precipitated by adding hexane as the solvent and acetone as the anti-solvent, followed by isolation via centrifugation [96]. The synthesis of Bi2Te3 via the sol–gel process involves heating a mixture of bismuth neodecanoate (0.63 mL), diphenyl ether (50 mL), and a thiol-capping agent (1-dodecanethiol, 10 mL) at 120 °C while trioctylphosphine telluride (1.5 mL) is injected. The solution is heated for 1 h followed by quenching in a cold water bath and isolation via centrifugation. The resulting black product is washed thoroughly with toluene and then dried in vacuum [97]. Similarly, BixSb2xTe3 is synthesized by heating a mixture of bismuth acetate (0.193 g), antimony acetate (0.342 g), diphenyl ether (50 mL), and a thiol capping agent (1-dodecanethiol,

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10 mL) at 120 °C while trioctylphosphine telluride (2 mL) is injected, using the same procedure as that for preparing Bi2Te3 [97]. 4.2.9. Microwave synthesis The synthesis of thermoelectric materials by microwaves is a rapid and reliable process. In a typical solid-state microwave synthesis method, a precursor (powder containing Pb, Se, or Te) is vacuum-sealed (105 mbar) in a quartz ampoule and exposed to microwave radiation in a microwave oven [98]. This technique can be used to synthesize nanocrystals of 15–100-nm size. Zhao et al. synthesized PbS NPs using precursors such as lead acetate trihydrate and thiourea with a weight ratio of 4:1 for Pb and S, respectively [99]. 4.2.10. Other methods Nakajima developed the method of ‘‘strip casting” for processing rare earth alloys such as the filled skutterudite-based alloy REx[Co1yMy]4Sb12 [100]. In this process, the raw materials are first heated to obtain a molten alloy, which is then quenched for solidification. The quenched alloy is further ground and sintered in order to obtain the final product. This process allows the simple and large-scale production of materials with a low production cost. Ota et al. showed that layered plate-shaped thermoelectric materials with a plastically deformed compact structure can be synthesized from raw materials [101]. Stefen and Ditscher introduced a faster method of synthesizing a direct thermoelectric element for a module by melting the material in an induction furnace [102]. In this technique, the melt is poured into a mold to obtain a module or sprayed onto a cold solid surface to obtain a pulverulent thermoelectric semiconducting material with particle size ranging from 1 lm to 5 mm [102]. Lee et al. developed another method to prepare anisotropically elongated thermoelectric materials: a multilayered nanowire structure is formed by the alternating deposition of thermoelectric layers and insulating layers via chemical vapor deposition (CVD), sputtering, or atomic-layer deposition [103]. Particles of Bi2Te3, Sb2Te3, Bi2TexSe3x, BiSb alloys, PbTe, skutterudites, and AgPbmSbTe2+m (LAST) are mixed with electrically or thermally insulating NPs or microparticles of oxides, nitrides, or fluorides. The particles are coated with an oxidizable metal (e.g., Al) via fluidized bed processing, evaporation, sputtering, CVD, or plasma spraying [103]. Murai and Kita developed a method of preparing nanocomposites with low thermal conductivities [104]. In this method, the resulting nanocomposite is composed of a matrix and a dispersed material in the form of NPs, the latter being either insulating or conducting. The surface of dispersed materials can be modified by organic molecules to obtain an efficient nanocomposite thermoelectric material. Murai and Kita developed a synthesis method for materials with high ZT values; here, the second phase is dispersed as NPs in a thermoelectric material matrix [105]. For example, Al2O3 particles are dispersed in a CoSb3 matrix [106]. Kuehling et al. proposed an extrusion method for synthesizing thermoelectric materials with Ar used as a protective gas [107]. In this process, the precursor material prepared by melt synthesis is cut into pieces ranging in size from 1 to 10 mm. This was followed by extrusion at 530 °C and 610 MPa to obtain a metallic shiny, compact, and cylindrically shaped material. This material is cut into slices of 1.5-mm thickness using a diamond saw. This process can be considered an alternative to the conventional melt-and-growth method. The variation in the thermoelectric properties of PbTe before and after the extrusion process is listed in Table 1. With high-pressure, high-temperature sintering, Malik synthesized PbTe with a high Seebeck coefficient within a short processing time of 15 min [108]. In this process, pressure and temperature play a key role in determining the thermoelectric properties of PbTe, as listed in Table 2. Ren et al. develop a method to enhance the thermoelectric properties of alloys: the alloys are subjected to one or more high-temperature annealing steps with the alloy exhibiting a mixed solid/liquid phase, followed by complete solidification during the cooling steps with the sample being cooled to solidify the melted regions, as shown in Fig. 13 [109]. Table 3 shows the effect of annealing temperature on the thermoelectric properties of the Si0.9Ge0.1 alloy. Other processing methods such as electrospinning, mechanosynthesis, and magnetron sputtering are also used to synthesize thermoelectric materials; the various factors affecting the performance of these materials are summarized in Table 4. 5. Strategies for enhancing the performance of thermoelectrics The increase in ZT with certain approaches/strategies can enhance the power factor, reduce thermal conductivity, or produce a combined effect of both factors. Degenerate semiconductors are most suitable for thermoelectric applications, as the

Table 1 Effect of extrusion process on the thermoelectric properties of PbTe [107]. Process

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K)

Thermal conductivity (W/m K)

ZT

Before extrusion

100 400

90 197

2700 650

21.87 25.22

2.7 1.45

0.3 1.17

After extrusion

100 300

90 170

2100 800

17.07 23.12

1.8 0.9

0.35 1.45

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C. Gayner, K.K. Kar / Progress in Materials Science 83 (2016) 330–382 Table 2 Effect of pressure, temperature, and processing time on the thermoelectric properties of PbTe [108]. Sample

Pressure (GPa)

Temperature (°C)

Time (min)

Powder size (mm)

Seebeck coefficient (lV/K) at 27 °C

As received A B C D E

N/A 6.5 6.5 6.5 6.5 6.5

N/A 900 1050 1050 900 900

N/A 5 15 5 15 5

N/A 2.0–4.0 2.0–4.0 2.0–4.0 <0.1 <0.1

232 235 357 225 309 340

Effect of pressure Synthesized F G H I

N/A 7.5 5.0 6.5 6.5

N/A 1200 1050 1050 1200

N/A 10 10 10 10

0.5 0.05–0.10 0.05–0.10 0.05–0.10 0.05–0.10

140 108 272 218 255

Table 3 Effect of annealing temperature on the thermoelectric properties of Si0.9Ge0.1 alloy. Si0.9Ge0.1 (n-type)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Thermal conductivity (W/m.K)

ZT (at 27 °C)

Unannealed Annealed at 1345 °C Annealed at 1365 °C

140 155 151

1052 768 845

9.73 7.67 6.00

0.063 0.072 0.096

Fig. 13. Schematic of (a) a few grains of an alloy; (b) after application of high-temperature annealing, which results in melting of peripheral layers of the grains; and (c) after solidification of the melted grain portions [109].

charge transport can be tuned by modifying their electronic band structure. Ideally, thermoelectric materials should possess charge carriers with optimized mobilities, carrier concentration, and effective mass. In this section, we briefly discuss various strategies such as doping, alloying, nanostructuring, use of nanocomposites/nanoinclusions, and use of thin films/superlattices to significantly enhance the performance of thermoelectric materials.

5.1. Doping When used, suitable doping elements can enhance the carrier concentration in a thermoelectric material by modifying their electronic band structure. Consequently, the net DOS are modified as the number of states per energy level is increased with further addition of the appropriate dopant atoms. The energy states close to the Fermi level within an order of kBT can promote electron transport. Hence, at an optimal carrier concentration, a higher Seebeck coefficient or ZT value can be obtained for semiconductors and semimetals. Therefore, a maximum ZT value can be obtained when the Fermi energy is closer to the conduction bands, when the degeneracy is not very strong. In such cases, E > EF and E < EF charge carriers are highly asymmetric in terms of their DOS and contribution [129,130]. With an increase in carrier concentration, the Fermi energy exceeds kBT, which has a negative effect on the Seebeck coefficient. In thermoelectric materials, elemental doping such as modulation or uniform doping is found to be essential for increasing the power factor. In particular, uniform doping tunes the carrier concentration, thereby enhancing the electrical properties while reducing the charge-carrier mobility. A heavily doped semiconductor exhibits a greater power factor than its undoped counterpart [13–17,131,132]. In PbTe and PbSe materials, some dopants distort the electronic DOS, increasing the Seebeck coefficient without affecting the electrical conductivity. Resonant states are observed in Tl:PbTe, Cr:PbTe, and Al:PbSe materials [13,17,53]. With the band engineering approach, modulation doping increases the mobility of charge carriers by reducing ionized impurity scattering. Here, the charge carriers are separated from the parent grains, in turn moving

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Table 4 Synthesis of thermoelectric materials via different routes and their thermoelectric properties. Material

Thermoelectric properties

Synthesis methods

Ref.

(BiSb)2Te3 + SiO2 and Bi2 (TeSe)3 + SiO2

a  100–250 lV/K (increases with Sb doping)

[110]

PbTe

Mechanically attributed mixture Ball milling or cryomilling

[112] [70]

Melt and growth

[113]

Direct reaction in a fused silica tube Mechanically attributed

[114]

Liquid-phase or solid-state precursor processing Mechanical grinding

[116]

(BiSb)2(TeSeS)3 + other dopants

ZT  0.20 for PbTe (micro-) ZT  0.20 for PbTe (nano-) ZT  0.29 for PbSe (micro-) ZT  0.36 for PbSe (nano-) at 27 °C 80% (2700 S/cm) at 47 °C and 25% (750 S/cm) higher electrical at 252 °C For Yb-doped sample, ZT  1.2 For Nd-, La-, and Ce-doped sample, ZT  0.95, 0.9, and 0.9, respectively At 127 °C, ZT  1.2, 1.05, and 0.9 for 1% NaSe0.5Te0.5, 1.5% TlTe and 1% NaTe addition, respectively Lowest thermal conductivity 0.55–0.7 W/m K ZT  1.5 at 500 °C In another sample, ZT  1.8 at 427 °C For p-type, ZT > 0.87 at 27 C For n-type, ZT > 0.8 at 27 °C Reduction in thermal conductivity Enhancement in ZT ZT  1.0 at 800 °C for p-type SiGe, ZT  1.3 at 900 °C for n-type SiGe ZT  1.45 at 110 °C for p-type BiSbTe alloy ZT is found to be increased

Chemical route (core–shelltype NPs) Solution containing a surfactant + Pb reagent + Te reagent

[118]

Fe–V–Al alloy + other elements

a  50–187 lV/K, r  1000–5000 S/cm, j  10–17 W/

Sintering of powder at normal pressure Casting process

Bi nanotubes cobalt oxide Pb2Sb6Te11  PbTe + Sb2Te3

m K at 27 °C. – – jlattice  1.5 W/m.K for 280–400-nm lamellar spacing.

Electrospinning process sol–gel process Melting a mixture of Pb, Te, Sb CVD

[120] [121]

Mechanosynthesis process

[123]

Prepared by mixing and heated in a tungsten crucible

[124]

Magnetron sputtering technique

[125]

Hydrothermal, functionalization process (core–shell structure) Ball-milled and hot-pressed

[126]

PbSe Bi–Sb matrix material + nanometal (Pb or Sn acetate) Filled CoSb3 (Filler atoms: rare earth elements)

Doped AgSbTe2

Tl–PbTe BiSbTeSe alloy BiTe + GaTe (or any other metal telluride) SiGe BiSbTe alloy

PbTe Bi2Te3

– – Thermopower in a range of 30–90 lV/K at 148 to 27 °C, respectively For Sr-filled clathrate, a  320 lV/K, r  78 S/cm at 27 °C For Eu-filled clathrate, a  150 lV/K, r  390 S/cm at 27 °C ZT  1.47–2.23 at 27 °C a  950 and 1250 lV/K for n- and p-type materials, respectively – –

Ca5ZnxAl2xSb6 Zintl compound PbTe–Ag2Te composite

ZT  0.8 at 727 °C jlattice  0.8–1.0 at 727 °C ZT  1.4 at 477 °C

Heterostructure of PbSe/PbSnSe

–

PbTe Bi2Te3 CoSb3 (+ alloying elements Ce, Fe, Ni) Clathrate + filler atoms such as Ba, Na, Ca, Sr, Eu, Ce, Ru

Si0.8Ge0.2/Si

Heating the 1st element and 2nd element Annealing the mixture Deposition

[111]

[115]

[117]

[119]

[122]

[127] [79]

[128]

into undoped grains [133]. In the carrier pocket engineering approach, doping causes the conducting electronic bands to converge in the bulk material with high valley degeneracy [15]. The convergence of the valence or conduction bands results in a simultaneous increase in both the Seebeck coefficient and the electrical conductivity [15]. The inclusion of anti-resonance NPs into the host material can be an alternative method of modulation doping or impurity doping for increasing ZT. Spherically symmetric core–shell particles of finite size, optimized effective mass, and band offset are used, which increase the power factor of a thermoelectric material by acting as resonant carrier scattering sites (termed as ‘‘invisible dopants”) [134]. Therefore, both the electrical conductivity and Seebeck coefficient are increased simultaneously in a host semiconductor embedded with anti-resonance NPs [134]. Further, a large acoustic mismatch is observed between the core–shell NP and the matrix material, which reduces the thermal conductivity without affecting the power factor [134,135]. In certain cases, energy levels of a dopant lying within a band of the host material may generate resonant energy levels, which can distort the DOS close to the Fermi level. Thus, the effective mass of the charge carriers can be increased without

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changing the carrier concentration in certain cases, for example, in Tl:PbTe, Al:PbSe, and Cr:PbTe. These resonant energy levels are crucial for increasing the Seebeck coefficient, which is dependent on the doping concentration. As shown in Fig. 14, significant changes in the DOS near the impurity band (if the Fermi level is positioned appropriately) can increase the thermopower of the material. This method increases the effective mass of the carriers without increasing the chargecarrier concentration, such that the electrical conductivity is maintained and only the Seebeck coefficient is increased [17,136]. Alternatively, if the bandgap is significantly greater than kBT (for non-degenerate charge carriers, in a parabolic band spectrum), then the concentration of minority carriers and their role in the charge transport become negligible. Here, the bandgap width and the Fermi level near the bottom of the conduction band are responsible for this unique behavior, thus significantly increasing the thermopower of the material. Therefore, for a given operating temperature, this spectral gap must be greater than the temperature, for instance, Eg > 10 kBT [80,137]. For PbTe, a maximum ZT value or Seebeck coefficient is achieved at 327 °C (at Eg  0.36 eV) [138]. Electronic bands are converged by tuning doping and composition, in turn leading to the convergence of many valleys in a bulk material. Large valley degeneracy is an ideal property for thermoelectric materials. Therefore, the separate pockets of a Fermi surface with the same energy (multiple degenerate valleys) can generate a large DOS effective mass without reducing the mobility. Here, the electronic performance depends on the weighted mobility l  (m⁄/me)3/2, where m⁄ is the DOS effective mass, l is the mobility of the charge carriers, and me is the mass of an electron [15]. In addition, the mobility is inversely 5=2

correlated with the band mass of a single valley or the mass of a single-pocket Fermi surface (l / 1=mb , where mb is the band mass) [15]. Hence, an increased band mass has a negative effect on the thermoelectric properties of a material. Therefore, the converged bands effectively lead to an increase in orbital degeneracy, Nv = (m⁄/mb)3/2, which can further increase m⁄. For example, the electronic bands in PbTe1xSex doped with 2% Na converge accompanied by an increase in the valley P of degeneracy of 16, as shown in Fig. 15. Here, the band itself contributes a valley of degeneracy of 12; hence, with this mechanism, a ZT value of 1.8 at 527 °C is obtained [15]. Modifications to the chemical composition of thermoelectric materials can be highly effective in increasing the thermopower while simultaneously reducing the lattice thermal conductivity. To achieve this, suitable alloying elements should induce mass fluctuation and improve the complex band structure. The disordered scattering of phonons leads to a decrease in thermal conductivity. This disorder is negligible in the case of electrons due to their larger velocities; thus, they have a considerably larger wavelength than phonons. Therefore, modification to the chemical composition is the primary method

Fig. 14. Schematic representation of the distorted band (the impurity distorts the DOS at the atomic level, where the bound state of impurity atoms distorts the band structure of host materials) [136]. In this figure a, b, c and d, temperature-dependent (a) electrical resistivity and (b) thermopower in the temperature range of 150–600 K. (c) At T = 300 K, the Fermi levels EFI, EFII, EFIII, and EFIV in samples (Cr:PbTe) I, II, III, and IV, respectively, do not align with the enhanced DOS. (d) Above 375 K, the Fermi level in samples I and II moves into the broadened region of enhanced DOS [17].

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Fig. 15. Schematic of relative energy of valence bands in PbTe0.85Se0.15 and temperature dependence of ZT of the PbTe1xSex materials doped with 2% Na. P Here, the L and valence bands appear to converge, which increases the valley of degeneracy of 16 [15].

of increasing the ZT and thermopower. However, mobility is an external factor that can affect the thermopower. All of these variations depend on the Mott relation [Eq. (6)]. For complex systems, the decrease in lattice thermal conductivity due to the disordered scattering of phonons via the difference in electronic states and the small mass fluctuation (2%) is found to enhance the thermoelectric performance [139]. In Tl9BiTe6, the marked changes in the thermoelectric performance are a result of the difference in electronic states (e.g., Tl1+ and Bi3+) [140]. In a homogeneous solid, the increase in thermopower is attributed to the energy-filtering effect. At the grain–grain interface, charge carriers encounter a potential barrier. The carrier that crosses this barrier shows a strong nonequilibrium energy distribution with high-energy particles, ultimately leading to an increased Seebeck coefficient [141]. The principle underlying this mechanism is that carriers with lower energies and higher Fermi energy cancel each other, thereby increasing the thermopower by eliminating low-energy carriers [141]. In addition, for a sintered sample, electrons are found to scatter across the potential barriers. These barriers are formed close to the grain boundaries, due to the localized states. These localized states are associated with point-like defects and grain-boundary dislocations, which can further increase the thermopower [139,142].

5.2. Alloying In a complex material system, alloying elements alter the performance of thermoelectric materials differently from the doping effect. Alloying increases the power factor while reducing the thermal conductivity. Ideally, such effects are observed in most HH alloys as well as in skutterudites and clathrates. The concept of PGEC is important particularly in the case of skutterudites. The guest atom in a matrix material acts as a scattering center, which reduces the thermal conductivity and enhances the electrical properties. The heavier guest atoms act as ‘‘rattling” atoms generating dynamic disorder, thus scattering phonons and creating a conducting path for the electrons [12,143–146]. In HH alloys, the electronic structure is modified by doping mechanisms. Alloying elements can introduce atomic disorder at an atomic site via an induced atomic fluctuation or the strain field effect, which reduces the lattice thermal conductivity [147,148].

5.3. Nanostructuring Nanostructuring is found to be effective in reducing the thermal conductivity by introducing nanoscale heterogeneities and nanodispersions. The quantum confinement effect and the energy-filtering effect are significant when the system size decreases and the length scale is comparable to its electron mean free path or wavelength. In this case, the DOS increase, eventually increasing the Seebeck coefficient. Conversely, the thermal conductivity is decreased by the scattering of phonons from the nanostructured surfaces or interfaces [30–33]. This approach is found to be effective in superlattices, nanowires, quantum dots, and thin-film systems [26–28]. In the case of bulk materials, for instance, PbTe, LAST, and LASTT, the nanoscale effects can be controlled by decreasing the grain size to nanometer regimes or using NPs. The presence of interfaces in nanostructured materials is the key parameter for reducing the lattice thermal conductivity [23–25,149]. Thus, using these interfaces can lead to a significant increase in ZT, particularly in the case of nanostructured materials. A nanocrystalline sample prepared by ball milling exhibits a ZT value of 1.4 at 127 °C in comparison with a BiSbTe alloy ingot [150]. Zhou et al. studied the effect of annealing on the thermoelectric properties, in relation to the microstructure of the system [23]. Based on their investigation, annealing increases the nanoscopic inhomogeneities within the grains (Fig. 16), in turn increasing phonon scattering and reducing the thermal conductivity, compared with the unannealed sample. However, nanostructuring has no effect on the power factor. After 30 days of annealing, a ZT value of 1.5 at 427 °C is obtained for Ag0.8Pb22.5SbTe20, which is 50% higher than that of the unannealed sample [23].

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347

Fig. 16. Microstructure of nanocrystalline BiSbTe alloys: (a) nanodots in matrix with small-angle grain boundaries and (b) nanodots in matrix without grain boundaries [150]; Ag0.8Pb22.5SbTe20 (c) unannealed and (d) annealed for 30 days [23]. Reproduced with permission from [23,150].

5.4. Nanocomposites and nanoinclusions Nanocomposite strategies that control phonon and electron transport are known to be effective in enhancing the thermoelectric performance of bulk materials with high thermal conductivity, such as HH alloys, skutterudites, and Si–Ge systems. In this study, the dispersed secondary NPs in the nanostructured host material are referred to as nanoinclusions. Strategies involving nanoinclusions are more effective due to the reduction in lattice thermal conductivity via phonon scattering, arising from the nanoinclusions, grain boundaries, and interfaces [42,73,139,151,152]. The concepts of spinodal decomposition as well as nucleation and growth are key to understanding ZT enhancement in nanocomposite materials [42,152–156]. Murai and Kita showed that ZT can be increased by dispersing second-phase NPs in a thermoelectric matrix, resulting in a larger mean free path for the charge carriers and a smaller mean free path for the phonons [105]. Here, thermal conductivity (j) decreases at a faster rate than the electrical conductivity. The addition of micro- or nanoinclusions is essential for increasing the thermopower because of the scattering of phonons or electrons at the hetero-boundaries. An increase in thermopower from 155 to 177 lV/K is observed with an increase in the ZrO2 particle content from 0 to 9 vol% [73]. The combined effect of nanostructuring and nanoinclusions leads to a significant increase in the Seebeck coefficient. The marked decrease in the effective hole density can be attributed mainly to the energy barrier (DE), which discriminates holes by trapping low-energy holes and only promoting high-energy holes via the full-Heusler (fH) valence band (Fig. 17). This mechanism leads to ‘‘hole culling,” and in turn an abrupt increase in mobility and effective mass of the holes. Therefore, the significant increase in thermopower is due to the simultaneous reduction in hole density and increase in the effective mass of holes. Here, an increase in mobility will not reduce the electrical conductivity [157]. In the case of bulk materials, nanoinclusions are favorable due to their low cost and easily processable materials. Nanocomposite materials are being increasingly used in thermoelectric devices due to their advantages, compared with bulk or isolated nanostructured materials [153–155,158]. Spinodal decomposition is used to create a material exhibiting compositional fluctuations at the nanoscale [152–156]. In spinodal decomposition, both compositions of the mixed-phase system share the same lattice, leading to a spatial modulation of the local composition at the nanoscale [152]. This spatial modulation is created coherently in the embedded NPs within the matrix, thereby forming nanostructured thermoelectric materials.

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Fig. 17. (a) Schematic of the atomic-scale band structure engineering of a p-type HH matrix using a nanoinclusion with HH composition, resulting in trapping of low-energy holes and transfer of high-energy holes at the HH heterojunction. (b) The decrease in effective carrier density and tuning of the carrier effective mass is seen [157].

Fig. 18. Structure of (PbTe)1x–(PbS)x system shows nanophase separation occurring via spinodal decomposition (stripes) and nucleation and growth (dot). Reproduced with permission from [152].

Matrix encapsulation along with nucleation and growth creates inclusions within the host matrix. During matrix encapsulation, the melted material is quenched immediately to introduce inclusions at the nanometer scale. The quenched samples can be annealed to improve their crystallinity as well as thermoelectric properties. Nanoscale inclusions can help reduce the thermal conductivity via phonon scattering and increase the electrical conductivity and Seebeck coefficient [158]. Fig. 18 illustrates the separation of the PbTe–PbS nanophases via spinodal decomposition and the nucleation and growth processes. Table 5 lists the synthesis routes and thermoelectric performances of certain nanocomposite thermoelectric materials. Biswas et al. proposed a hierarchical architectural approach in a p-type PbTe–PbSr system, to obtain a ZT value of 2.2 at 642 °C [18]. This approach is based on the nanostructuring effect arising from phonon scattering at multiple length scales. As shown in Fig. 19, the scattering processes at different length scales can be attributed to atomic-scale alloy scattering, mesoscale grain-boundary scattering, and scattering from nanoscale endotaxial precipitation. These multiscale hierarchical structures exhibit greater ZT values than their individual counterparts. When combined, atomic-scale alloy doping and endotaxial nanostructuring with a mesoscale grain boundary can significantly increase the ZT value compared to the nanostructuring effect alone [18].

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C. Gayner, K.K. Kar / Progress in Materials Science 83 (2016) 330–382 Table 5 Synthesis routes and thermoelectric properties of nanocomposite thermoelectric materials. Material

Thermoelectric properties

Synthesis route

Ref.

(Co, Ni)Sb3/30 vol% CeO2 (Bi, Sb)2Te3/30 vol% Te

Liquid-phase synthesis

[104]

CoSb3 + Sb2Te3

ZT  1.24 (a  203 lV/K, j  0.7 W/m.K) at 400 °C ZT  1.15 (a  203 lV/K, j  0.45 W/ m.K) at 27 °C ZT  0.6 at 327 °C

[105]

CoSb3 + Al2O3

–

PbTe + SiO2 Bi2Te3 + SiO2 PbSexTe1x/PbTe quantum dot Zn4Sb3-based alloys and ZrNiSn-based HH alloys. SiGe, PbSe, PbTe + nanoinclusions

–

Solution or slurry containing elements + reducing agent Dispersion of elements in liquid containing ceramic particles followed by hydrothermal process Hydrothermal, functionalization process (core–shell structure) Vapor deposition, electrochemical deposition, and combination of the two. Wet chemistry process

Bi0.5Sb1.5Te3 + Ag acetate/Cu acetate Si or PbSe NP-embedded to Ge or PbTe PbTe – Sb, Bi, InSb, CdTe, PbS

jlattice  1.38 W/m.K at 27 °C jlattice  2 W/m.K at 27 °C Power factor of nanocomposite 2– 100 lW/cm K2 ZT of 1.35 at 127 °C (j  0.75 W/m K, power factor 26 lW/cm K2) – For PbTe–CdTe, power factor 30 lW/cm K2 at 27 °C For PbTe–PbS, power factor 28 lW/ cm K2 at 27 °C

[106] [126] [153] [154]

Ball milling

[155]

Wet chemistry or vapor deposition process Spinodal decomposition, melting, quenching, and annealing for a long period

[156] [158]

Fig. 19. ZT values for the respective length scales: the atomic scale (alloy scattering: red, Te; blue, Pb; green, dopant), the nanoscale (PbTe matrix, gray; SrTe nanocrystals, blue) to the mesoscale (grain-boundary scattering). Reproduced with permission from [18].

5.5. Thin films and superlattices Superlattices are periodic layers of two or more materials with each layer of nanometer thickness. Synthesizing superlattices with a controlled microstructure is cumbersome. Two types of superlattice structures are seen: the ‘‘homostructure” corresponding to a single matrix material and the ‘‘heterostructure” corresponding to two or three different materials. Superlattices are synthesized via molecular beam epitaxy [30,32], sputtering [33], metal organic CVD [34,35], etc. Ultrathin-film nanostructures are synthesized by thermal evaporation [36], electrodeposition [37], pulsed laser deposition [29], CVD [38], etc. Harman et al. synthesized a quantum dot superlattice (QDSL) composed of PbTe/PbSe0.98Te0.02 doped with Bi on a BaF2 substrate with ZT values of 1.6 and 3.5 at 300 and 570 K, respectively [30]. Venkatasubramaniam et al. reported the synthesis of multiple-quantum-well Bi2Te3/Sb2Te3 superlattices with a ZT value of 2.4. High ZT values can be attained by reducing the thermal conductivity in the order of 0.4–0.6 W/m K by controlling the transport of phonons and electrons without any effect on the power factor [31]. The thermoelectric properties of various superlattices and thin films are listed in Table 6. The major drawbacks of using superlattices are the complex synthesis routes, high cost, and inability to support a large temperature difference across the material [31–41]. 6. Performance evaluation of thermoelectric materials 6.1. Effect of material properties on performance The performance of thermoelectric materials is strongly affected by the crystal structure, microstructure, densification, and grain orientation and size. The presence of secondary phases, such as metallic phases, can reduce the performance of complex material systems, whereas an increase in temperature and pressure may lead to grain orientation. Therefore, these material parameters can be used to control the thermoelectric properties.

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Table 6 Thermoelectric properties of various superlattices and thin films. Material

Operating temperature (K)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm.K2)

Thermal conductivity (W/m.K)

ZT

Ref.

Sb Bi Te p-Bi0.5Sb1.5Te3 (Deposited on mica) p-Bi0.5Sb1.5Te3 (Deposited on AIN/Si) n-(Bi2Te3)90(Sb2Te3)5 (Sb2Se3)5 (Deposited on mica) n-(Bi2Te3)90(Sb2Te3)5 (Sb2Se3)5 (Deposited on AIN/Si) Bi-doped (n-type) QDSL-A PbSe0.98Te0.02/PbTe Bi-doped (n-type) QDSL-B PbSe0.98Te0.02/PbTe n-(BiSb)2 (SeTe)3 p-type Bi2Te3/Sb2Te3 n-type Bi2Te3/Bi2Te2.83Se0.17

300 300 300 300 300 300

47 70 500 86 142 53

0.0288  106 0.00867  106 2 96 52 584

63.61 42.48 0.5 0.71 1.04 1.64

24.3 7.87 2.35 0.3 0.4 0.3

0.078 0.161 0.63  104 0.05 0.08 0.17

[159] [159] [159] [29] [29] [29]

300

155

105

2.52

0.3

0.25

[29]

300 300 300 300 300

219 208 228 210 238

728 584 806 1054 813

34.91 25.26 41.89 46 46.06

0.58 0.58 1.36 0.58 0.945

1.6 1.3 0.9 2.4 1.46

[30] [30] [30] [31] [31]

6.1.1. Crystal structure For complex thermoelectric materials, phase analysis is needed to determine their contribution. The presence of secondary phases can affect the thermoelectric properties; for example, each metallic or semiconducting phase has a different crystal structure determining the material properties [54,88]. In the Cu2Se system, the low-temperature a phase crystallizes into a monoclinic structure, which is stable up to 127 °C. At elevated temperatures, however, it can turn into a cubic structure [88]. By contrast, the high-temperature b phase has a face-centered cubic (FCC) structure with the lattice parameters a = b = c = 5.8639 Å at 200 °C. The lattice parameters reported for the monoclinic a phase are a = b = 11.52 Å and c = 11.74 Å. At room temperature, the (2 2 2) plane of the a phase is transformed into the (1 1 1) plane of the b phase at 200 °C as a result of the phase transformation and disorder at the Cu sites [54]. The effects of phases on the thermoelectric properties of materials are plotted in Fig. 20. 6.1.2. Microstructure/Densification The microstructure of a material is key to determining its thermoelectric performance. Nanostructured materials exhibit low lattice thermal conductivities and high ZT values, compared with bulk materials. The ZT value can be increased via microstructural engineering, which helps maintain a larger mean free path for the charge carriers and a smaller mean free path for the phonons. The difference in the mean free path between phonons and charge carriers can be used to obtain higher ZT values, as the thermal conductivity (j) decreases at a faster rate than electrical conductivity does. The mean free paths of the carriers and phonons are calculated using Eqs. (9) and (10), respectively [105]:

lc ¼

ðlm  v Þ e

lp ¼

3jLattice CP v a

ð9Þ ð10Þ

where lc is the mean free path of carriers, lp is the mean free path of phonons, l is the mobility, m⁄ is the effective mass of charge carriers, v is the carrier velocity, e is the electric charge, jLattice is the lattice thermal conductivity, Cp is the specific heat, and va is the acoustic velocity. The thermoelectric properties of bulk materials can be altered by proper tuning of their microstructure. For instance, nanoscopic inhomogeneities, such as an increased number of grain boundaries to the parent system, are found to enhance the thermoelectric performance of SiGe and BiSbTe alloy [150,160]. Material with nanoscopic inhomogeneities exhibit lower lattice thermal conductivities. Techniques such as SPS and HP are used to reform the microstructure and densification of materials. For example, the microstructure of HfNiSn0.98Sb0.02 synthesized by different methods such as levitation melting, SPS, and HP differ, as shown in Fig. 21. HfNiSn0.98Sb0.02 synthesized by SPS exhibits a greater electrical conductivity than that prepared by HP [86]. 6.1.3. Grain orientation and size In polycrystalline bulk materials, the grain orientation and interfacial microstructure are closely associated with each other. These parameters are crucial for determining the thermoelectric performance of a material. If the preferential alignment of grains is not in the direction of electron transport, as shown in Fig. 22a, the electrical properties of the material are diminished. The preferential alignment of grains in the direction of electron transport, as shown in Fig. 22b, can be achieved with processes such as annealing or sintering. A reduction in grain size produces more grain boundaries, which in turn scatter the phonons more effectively, thereby reducing thermal conduction via the interfacial scattering process, as shown in Fig. 22c. In the case of nanostructured materials, controlling the grain size is crucial not only for single-phase materials but also for multiphase systems, such as materials with dispersed secondary phases. Nanocomposites or nanoinclusions

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351

Fig. 20. Effect of a and b phases and temperature on (a) electrical resistivity (open symbols) and Seebeck coefficient (closed symbols) and (b) thermal conductivity of Cu2Se. Reproduced with permission from [88].

can be used to provide a path for the interfacial scattering and scattering from the nanoinclusions, as shown in Fig. 22d–f [161]. Nanostructuring is found to be an effective method of reducing the thermal conductivity via the strong phonon scattering from their interfaces [19,23,24,150]. Heremans et al. studied the effect of grain size on the Seebeck coefficient [162]. Thermoelectric nanogranular materials are known to exhibit narrow grain-size distribution. A range of grain sizes can be obtained, according to Eq. (11):

lc < d < 5lc 2

ð11Þ

where d is the grain size and lc is the mean free path of charge carriers. Here, the mean free path is the phonon-limited mean free path of electrons or holes [162]. According to Eq. (11), the mean free path of a moving particle is dependent on size dependent: it decreases for smaller grain size and increases for larger grain size. For smaller grain size, the grain boundaries are effectively larger, which scatter the charge carriers and phonons more readily; thus, the mean free path is affected by the grain size. 6.2. Thermoelectric properties of inorganic materials Inorganic materials have been traditionally used in thermoelectric devices, as they are superior to organic materials in their thermoelectric properties. Inorganic materials are classified into various groups such as intermetallics, skutterudites, clathrates, HH, oxides, rare earth chalcogenides, Zintl-phase materials, pnicogens, and nitrides. The crystal structure, physical properties, and thermoelectric performance of inorganic materials are discussed in this section. 6.2.1. Skutterudites The term ‘‘skutterudite” refers to naturally occurring mineral CoAs3, first discovered in Skutterud, Norway. The cubic structure of skutterudites is composed of 32 atoms per unit cell [43]. CoAs3 exhibits a distorted version of the AB3-type

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Fig. 21. SEM images of (a) levitation-melted HfNiSn0.98Sb0.02, (b) levitation-melted, SPS-processed HfNiSn0.98Sb0.02, and (c) levitation-melted, hot-pressed HfNiSn0.98Sb0.02. Reproduced with permission from [86].

Fig. 22. Strategies for improving the bulk thermoelectric performance involve several distinct classes of grain and interfacial microstructure: (a) polycrystalline microstructure, (b) preferential alignment of grains along the transport direction, (c) and reduced grain size for favorable interfacial scattering process; and (d–f) nanocomposites: (d) nanocoated grains, (e) nanoinclusions, and (f) lamellar/multilayer structure. Reproduced with permission from [161].

perovskite structure, that is, MX3 (M = Co, Rh, Ir; X = P, As, Sb), with an octahedral structure and a void at its center. The voids are large enough to host large metal atoms to form filled skutterudites. The void-filling atom can act as an electron donor or electron acceptor, thus changing the electron concentration. Moreover, these void-filling atoms act as strong phononscattering centers, which can reduce the lattice thermal conductivity. Smaller and heavier atoms in the voids introduce significant disorder to the lattice, leading to a reduction in the lattice thermal conductivity [43,44]. In the case of CoSb3, the filler atoms inside the voids donate their valence electrons to the CoSb3 framework, eventually transforming the material into a heavily doped semiconductor. In CoSb3, 18 valence electrons are present, nine of which are shared by Co and three are shared by each Sb atom. It is a complex phase system with only the d-phase exhibiting semiconducting properties, which are stable up to 874 °C. Thus, it is suitable for high-temperature thermoelectric applications. It is an indirect-band semiconductor with a bandgap of 0.57 eV. The Co-rich CoSb3 phase acts as an n-type thermoelement, whereas the Sb-rich phase acts as a p-type thermoelement. The properties of the CoSb3 system can be tuned by substituting foreign atoms in place of Co and Sb sites, such as substituting the Co sites with Fe or Ni and the Sb sites with Sn or Te. The substituted atoms may act as active dopants by donating (e.g., Ni and Te) or accepting (e.g., Sn) electrons [143,144]. Table 7 summarizes the properties of some elements used as dopants in CoSb3. CoSb3 exhibits a power factor of 6.5 lW/cm K. However, it displays higher thermal conductivities of 10 W/m K at 27 °C and 4–5 W/m K at 427 °C [20,145] compared with Te-based materials, such as PbTe and Bi2Te3. At this high thermal conductivity, high ZT values cannot be obtained. For the pristine (undoped) CoSb3 system, a ZT value of 0.1 is obtained at 323 °C. Studies have shown that filler atoms inside the CoSb3 matrix can enhance the thermoelectric response of the doped CoSb3 system [145]. For the YbxCo4Sb12 system, ZT values of 0.3 and 1 are obtained at 27 and 327 °C, respectively. With a Yb filler, the ZT is increased by a factor of 7 and 10, respectively, compared with an undoped CoSb3 system at the same temperatures [89,145,163,164]. When guest atoms form weak bonds with the host lattice, phonon transport occurs more readily. The ‘‘rattling effect” observed in PGEC-type materials reduces the lattice thermal conductivity without affecting their

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C. Gayner, K.K. Kar / Progress in Materials Science 83 (2016) 330–382 Table 7 Physical and thermoelectric properties of various dopants used in the CoSb3 system [159]. Dopant

Seebeck coefficient (lV/K) at 27 °C

Seebeck coefficient (lV/K) at 227 °C

Electrical conductivity (S/cm)  106

Thermal conductivity (W/m K)

Melting point (°C)

Crystal structure

Atomic radius (Å)

Co Sb Yb Ca Ba Ce La

30.8 47 30 10.3 12.1 6.2 1.7

40.8 – 20.2 17.1 28.5 5.2 2

0.172 0.0288 0.0351 0.298 0.03 0.0115 0.0126

100 24.3 34.9 201 18.4 11.4 13.5

1495 630.9 824 839 729 798 920

HCP Rhombohedral FCC FCC BCC FCC Hexagonal

1.67 1.53 2.4 2.23 2.78 2.7 2.74

electrical conductivity [12]. Recent studies have indicated that a reduction in the grain size or nanostructuring can enhance the thermoelectric properties of a CoSb3 system to a greater extent [20,21]. A thermal conductivity of 1.61 W/m K is obtained for bulk CoSb3 prepared from hydrothermally synthesized nanopowder using SPS or HP [20–22]. An efficient thermoelectric device must possess n- and p-type legs with similar mechanical and thermal properties so as to minimize failure due to thermal stresses. In this context, CoSb3 is a suitable candidate for use in thermoelectric devices operating at medium temperatures due to its good mechanical and thermal properties [4,20,145]. PGEC materials exhibit a regular periodic cage-like structure, in which the electron can move freely as in an ideal crystal. A large, heavier, and weakly bonded guest atom (or ‘‘rattler”) creates a disturbance in the oscillations due to lattice vibrations. This guest atom can reduce the heat conduction in the material, resulting in a glass-like low thermal conductivity [164]. Shi et al. reported a reduction in the lattice thermal conductivity by filling Ba and Yb atoms in the voids of a skutterudite structure, which afforded a broad range of resonant phonon scattering [165]. A ZT value of 1.36 is obtained for n-type BaxYbyCo4Sb12 at 527 °C. The multiple-filled skutterudite Ba0.08La0.05Yb0.04Co4Sb12 exhibits a ZT of 1.7 at 577 °C. Multiple filling of filler atoms is more efficient in reducing the lattice thermal conductivity, as in the case of Ba0.08La0.05Yb0.04Co4Sb12 [165]. The high ZT value can be attributed to the SPS process, as it improves the microstructure of the material and ensures homogeneous multiple filling of atoms in the voids. The effect of filler atoms on the thermoelectric properties of skutterudites at different temperatures are summarized in Table 8. Table 9 lists the thermoelectric properties of certain skutterudite materials. Table 8 Effect of filler atoms on the thermoelectric properties of skutterudites at different temperatures. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/mK)

ZT

Ref.

CoSb3 YbxCo4Sb12(x  0.19) Ba0.1Yb0.2Co4Sb12 Ba0.1Yb0.2In0.1Co4Sb12 Ba0.1Yb0.2In0.2Co4Sb12 Ba0.1Yb0.2In0.25Co4Sb12 Ba0.15Yb0.2In0.2Co4Sb12 CoSb3 YbxCo4Sb12 (x  0.19) Ba0.1Yb0.2Co4Sb12 Ba0.1Yb0.2In0.1Co4Sb12 Ba0.1Yb0.2In0.2Co4Sb12 Ba0.1Yb0.2In0.25Co4Sb12 Ba0.15Yb0.2In0.2Co4Sb12 Ba0.06La0.05Yb0.02Co4Sb12 Ba0.08La0.05Yb0.04Co4Sb12 Ba0.10La0.05Yb0.07Co4Sb12 Ba0.09La0.04Yb0.13Co4Sb12 Ba0.09La0.04Yb0.14Co4Sb12 Ba0.06La0.05Yb0.02Co4Sb12 Ba0.08La0.05Yb0.04Co4Sb12 Ba0.10La0.05Yb0.07Co4Sb12 Ba0.09La0.04Yb0.13Co4Sb12 Ba0.09La0.04Yb0.14Co4Sb12 In0.2Yb0.2Co4Sb12 Ce0.1In0.15Yb0.15Co4Sb12 Ce0.1In0.1Yb0.2Co4Sb12 In0.2Yb0.2Co4Sb12 Ce0.1In0.15Yb0.15Co4Sb12 Ce0.1In0.1Yb0.2Co4Sb12

27 27 27 27 27 27 27 327 327 527 527 527 527 527 27 27 27 27 27 577 577 577 577 577 27 27 27 527 527 527

120 150 138 137 126 123 117 180 220 213 205 178 209 185 138 126 107 104 93 210 192 185 174 161 141 143 174 216 208 225

400 2000 1750 1779 1950 2065 2345 200 1110 960 1040 1200 1280 1410 1831 2398 3003 3058 3367 1052 1428 1666 1818 2220 1500 1500 1500 1100 1200 1000

5.7 45 32 33 31 31 32 6.5 53.4 43.5 43.7 38.2 55.9 48.2 35 38 34 33 29 45.48 52.6 57 55.04 57.54 29.82 30.67 45.41 51.32 51.91 50.62

10 5–6 4.03 3.76 3.60 3.55 3.55 6 3.35 4.08 3.5 3.5 3.38 3.32 3.0 2.7 3.1 3.0 3.0 3.0 2.6 3.1 3.3 3.47 3.6 3.47 3.34 3.35 3.04 3.08

0.017 0.3 0.23 0.26 0.25 0.26 0.27 0.1 1.08 0.85 0.99 0.87 1.32 1.16 0.35 0.42 0.33 0.33 0.29 1.28 1.71 1.56 1.41 1.40 0.25 0.26 0.40 1.22 1.4 1.34

[145] [145] [163] [163] [163] [163] [163] [145] [145] [163] [163] [163] [163] [163] [89] [89] [89] [89] [89] [89] [89] [89] [89] [89] [164] [164] [164] [164] [164] [164]

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Table 9 Thermoelectric properties of skutterudite materials [166]. Material

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K)

Thermal conductivity (W/m K)

ZT at 327 °C

In0.2Ce0.1Co4Sb12 In0.2Ce0.2Co4Sb12 In0.25Ce0.1Co4Sb12 In0.25Ce0.2Co4Sb12 In0.2Ce0.1Yb0.1Co4Sb12 In0.2Ce0.1Sm0.1Co4Sb12 In0.2Ce0.1Tb0.1Co4Sb12 In0.2Ce0.1Ho0.1Co4Sb12

340 295 325 300 325 305 348 310

230 330 500 450 320 280 280 230

26.58 28.46 52.8 40.5 33.8 26 33 22.1

1 1 2.2 1.6 1.65 1.15 1.4 1

1.6 1.7 1.4 1.5 1.2 1.3 1.4 1.3

Yang et al. introduced a filled-skutterudite system with the formula GyM4X12, where G denotes a rare earth element; M denotes Co, Rh, or Ir; and X denotes Sb, P, or As [56]. Multiple-filled skutterudite systems, such as Ba0.08Yb0.04La0.05Co4Sb12, exhibit a ZT value of 1.8 at 900 K, the highest value reported for skutterudites [56]. Fig. 23 compares the ZT values of filled skutterudites with the commercially available materials at different temperature regimes.

6.2.2. Telluride-based materials: PbTe PbTe and Bi2Te3 are conventional state-of-the-art thermoelectric materials that have been used for the past 50 years. Te-based materials exhibit low thermal conductivity (2.3 W/m K for PbTe and 1.7 W/m K for Bi2Te3) and a high Seebeck coefficient of 500 lV/K at room temperature. PbTe is a promising thermoelectric material that can efficiently operate at medium temperatures (327–527 °C). It crystallizes into the isomorphous cubic NaCl crystal structure with Pb atoms at the cationic sites and Te at the anionic sites; the Pb atom is located at the origin and the Te atom is at (1/2, 1/2, 1/2). Due to its bandgap (0.32 eV), PbTe can be doped either with n-type or with p-type dopants. Pb-rich PbTe is an n-type semiconductor, whereas Te-rich PbTe is a p-type semiconductor. The dopant atoms can introduce electronic states in PbTe, thus distorting the electronic DOS and increasing the Seebeck coefficient (290 lV/K in Tl:PbTe and 220 lV/K in undoped PbTe at 427 °C) [13]. A resonant state and pinning of the energy levels are noted in Tl- and Cr-doped PbTe, resulting in their increased ZT values [13,17]. Strategies such as nanostructuring can be used to obtain a high Seebeck coefficient in PbTe [22]. The thermal conductivity and electronic properties of PbTe are determined by its microstructure due to grain and grain-boundary refinement [131]. The increase in DOS due to doping can be visualized via low-temperature-specific heat and optical spectroscopy analyses. Simultaneous doping of more than one element in a thermoelectric material can yield a high power factor. For instance, a high power factor of 47 lW/cm K2 is obtained for Pb0.91Cr0.009Te co-doped with I [17,90]. A thermal conductivity of 1.2 W/m K and a ZT value of 2.2 at 277 °C are reported for this material [17]. Here, Cr-doped PbTe causes a distortion in the conduction band, locally enhancing the DOS, whereas I tunes the position of the Fermi level to control the electron density. Recently, various approaches such as resonant impurities, band convergence, and alloying have been proposed for enhancing the thermoelectric properties of PbTe. Wang et al. studied the effect of processing conditions on the performance of p-type 2% Na-doped PbTe by varying the HP pressure from 70 to 130 MPa and the sintering time from 0.5 to 2 h. The microstructure of the material can be controlled by varying the sintering time and pressure. For a p-type 2% Na-doped PbTe, a ZT of value of 1.74 is obtained at 501 °C [15]. This increase in ZT can be attributed to the electronic

Fig. 23. Effect of temperature on the ZT of filled skutterudites and commercially available thermoelectric materials [56].

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and lattice contributions to the thermal conductivity [167]. Kanatzidis et al. proposed a hierarchical architectural approach to obtain a maximum ZT value of 2.2 at 642 °C in a p-type PbTe–PbSr system [18]. Beyond nanostructuring, this approach indicates various effects at multiple length scale scattering phonons, including atomic-scale alloy scattering, scattering from mesoscale grain boundaries, and scattering from nanoscale endotaxial precipitation [18]. This panoscopic approach is a combination of band engineering and nanostructuring, particularly for optimizing the ZT values of thermoelectric materials [18]. Pei et al. proposed a carrier pocket engineering approach yielding a ZT of 1.8 for a 2% Na-doped PbTe0.85Se0.15 bulk sample at 527 °C [15]. Heavy doping in the PbTe1xSex system results in high valley degeneracy via the convergence of the conduction band [15]. The effects of various dopants on the thermoelectric properties of the PbTe system are summarized in Table 10. A nanocomposite thermoelectric material was developed using a compound with the formula MQ (M = Ge/Sn/Pb and Q = S/Se/Te), Ge(1x)Six, and ZnTe, and its thermoelectric properties were evaluated [168]. A Vickers hardness P0.4 GPa was obtained for the nanocomposite prepared by mixing PbTe with Ge(1x)Six, when compared to that of pristine PbTe. A high ZT value of 1.2 was reported for PbI2-doped PbTe at 477 °C, which can be attributed to the increased electrical conductivity of the complex system [168]. A reduction in lattice thermal conductivity was observed in PbTe-Ge(1x)Six with varying composition, as shown in Table 11. Kanatzidis et al. proposed a complex system with the formula MQ–AB composed of dispersed nanoscale inclusions with a rock-salt structure [169]. The nanoinclusions can inhibit heat flow via the strong acoustic phonon scattering at the matrix– nanoinclusion interfaces without decreasing the carrier mobility. However, a consequent decrease in the lattice thermal conductivity was observed. Table 12 summarizes the thermoelectric properties of MQ–AB-type complex material systems.

Table 10 Effects of various dopants on the thermoelectric properties of PbTe. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Pb Te Tl Na Se Ce PbTe PbTe:Na (2%) Tl0.02Pb0.98Te 2% Na-doped Na0.02Mg0.03Pb0.95Te 2% Na-doped PbTe0.75Se0.25 Pb0.96Mn0.04Te:Na Pb0.95Ce0.05Te Pb0.91Cr0.009Te co-doped with I PbTe (SPS) PbTe:Na (2%) Tl0.02Pb0.98Te 2% Na-doped Na0.02Mg0.03Pb0.95Te 2% Na-doped PbTe0.85Se0.15 2% Na-doped PbTe0.85Se0.15 Pb0.96Mn0.04Te:Na Pb0.95Ce0.05Te Pb0.91Cr0.009Te co-doped with I

27 27 27 27 27 27 27 27 27 27

1.05 500 – 2 900 6.2 225 62 130 110

0.0481  106 2 0.0617  106 0.21  106 1  106 0.0115  106 333 2500 520 1000

0.0530 0.5 0.84 8.1  107 0.442 16.85 9.61 8.78 12.1

35.3 2.35 46.1 141 2.04 11.4 1.45 3.7 1.8 2.6

0.45  104 0.63  104 – 0.17  105 1.19  108 1.16  103 0.34 0.077 0.146 0.140

[159] [159] [159] [159] [159] [159] [131] [15] [13] [14]

27

56

1612

5.05

2.7

0.056

[15]

27 27 27

100 150 100

1333 357 1666

14 8.03 16.66

2.2 1.0 If 1.8

[16] [131] [17]

427 427 527 427

310 225 325 294

105 543 172 232

10.09 27.48 18.16 20.05

1.1 1.5 1.0 1.3

0.2 0.24 0.27 (expected) 0.65 1.28 1.5 1.07

527

220

526

27.82

1.2

1.8

[15]

427

202

590

24.07

1.3

1.3

[15]

427 427 277

263 220 275

333 208 625

23 10.06 47.26

1.0 0.8 If, 1.2

1.6 0.88 2.2 (expected)

[16] [131] [17]

[131] [15] [13] [14]

Table 11 Lattice thermal conductivity variation of PbTe-Ge(1x)Six with composition [168]. Material

Lattice thermal conductivity (W/m K) at 327 °C

PbTe–Ge (20%) PbTe–Ge0.95Si0.05 (20%) PbTe–Ge0.80Si0.20 (20%) PbTe–Ge0.70Si0.30 (20%)

1.50 1.55 1.20 1.30

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Table 12 Thermoelectric properties of MQ–AB-type complex material systems [169]. Material

Temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

1% Na2Te–Sr0.01PbTe1.01 1% Na2Te–Sr0.02PbTe1.02 1% Na2Te–Sr0.03PbTe1.03 PbTe–SrTe (2%) PbTe–PbSnS2 (6%) PbTe–PbSnS2 (3%) + 0.055 mol% PbI2 Na-doped PbTe–5 mol% CaTe

427 427 427 527 427 427 527

275 280 282 285 285 240 265

270 250 260 250 90 320 350

20.41 19.60 20.67 20.30 7.31 18.43 24.57

1.15 1.16 1.16 0.95 0.78 0.90 1.15

1.24 1.20 1.24 1.70 0.65 1.43 1.70

Table 13 Thermoelectric properties of doped PbTe material systems at 27 °C [170]. Material

Seebeck coefficient (lV/K) (Tcold = 50 °C, Thot = 280 °C)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Pb0.992Ge0.005Ti0.003Te1.003 Pb0.992Ge0.005Zr0.003Te1.003 Pb0.99Bi0.005Zr0.005Te1.001 Pb0.989Ge0.01Ag0.001Te1.001 Pb0.987Ge0.01Sn0.03Te1.001 Pb0.997Zr0.03Te1.003 Pb0.999Ag0.01Te1.003 Pb0.995Cu0.05Te1.003

165.4 132.1 154.6 326.5 290.4 139.4 314.5 136.7

1641.6 2485.9 992 407.3 249.4 3895.7 451.2 1936.5

44.9 43.4 23.7 43.4 21.0 75.7 44.6 36.2

These nanoinclusions can increase ZT values up to 1.7 at 527 °C, lowering the thermal conductivity without affecting its electrical conductivity [169]. Haass developed novel systems comprising doped PbTe with the general formula Pb(1X1+X2+X3) A1x1A2x2 . . .. AnxnTe1+z; these systems displayed high electrical conductivity as well as high power factor [170]. Table 13 summarizes the thermoelectric properties of these novel doped PbTe complex material systems. The doping process was found to increase the power factor of the material system, with a high value of 75.7 lW/cm K2 obtained for the Zr-doped PbTe material system, because of high electrical conductivity [170]. Sengupta et al. synthesized PbSe-coated PbTe NPs via a sonochemical method, which exhibited enhanced thermoelectric properties [39]. PbTe exhibits a thermal conductivity of 1.16 W/m K in its uncoated pristine form and a low value of 0.8 W/ m K when surface-coated with PbSe. This reduction in thermal conductivity results in an increased scattering of phonons at the PbTe/PbSe interfaces. The thermopower of the PbSe-coated PbTe NPs is increased up to 550 lV/K, which is 5–10 times higher than that of the uncoated pristine PbTe. This increase in thermopower can be attributed to their small particle size and the presence of interfaces [39]. 6.2.3. Bi2Te3 Bi2Te3 is a low-temperature thermoelectric material (can operate up to 227 °C) first proposed in 1954 [171]. It crystallizes into a hexahedral-layered structure with five atomic layers (Te1–Bi–Te2–Bi–Te1) stacked by van der Waals interactions along the c-axis of the unit cell [171]. The Te2 atoms are covalently bonded to the surrounding six Bi atoms by sp3d2 hybrids. Bi2Te3 displays unique properties such as a high Seebeck coefficient (220 lV/K), good electrical conductivity (400 S/cm), and low thermal conductivity (1.5 W/m K). It has a low melting temperature of 585 °C. At room temperature, the bandgap of Bi2Te3 is 0.13 eV [171]. The effects of composition and processing routes on the thermoelectric properties of Bi2Te3-based material systems are summarized in Table 14. Tang et al. prepared Bi2Te3 bulk material with a layered structure by the melt-spinning and SPS processes [172]. During this synthesis, they found an increase in electrical conductivity and decrease in thermal conductivity when the rolling speed was increased. A maximum ZT of 1.35 was obtained at room temperature [172]. Zhao et al. prepared nanotubes of quasilayered Bi2Te3 via a hydrothermal method, the low dimensionality and hollow structure of which increase ZT value via phonon-blocking mechanisms [87]. Poudel et al. synthesized an alloy composed of Bi and Sb and investigated the effect of alloying on its thermoelectric properties [150]. The alloying process yields a ZT of 1.4 for the Bi–Sb alloy [150]. Xie et al. reported a nanostructuring approach for synthesizing a p-type Bi2Te3 system [57]. This system displays a low lattice thermal conductivity with no reduction in its electrical conductivity (a  225 lV/K, r  625 S/cm, and jlattice  0.5 W/ m K) and a high ZT of 1.5 at 117 °C [57]. With the nanostructuring procedure, the ZT was increased by 50% compared with that achieved with commercially available zone-melted materials [57]. A p-type Bi0.5Sb1.5Te3 nanograin architecture with grain size of 20–50 nm yielded a ZT of 1.36 at 87 °C [173]. This is due to the significant decrease in thermal conductivity (0.9 W/m K). High-density nanostructured Bi2Te3 pellets were synthesized by a wet-chemical method followed by CP and field-assisted sintering, displaying a ZT value of 0.38 at 27 °C [174]. A ZT of 0.95 was obtained at 27 °C for chemically

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Table 14 Effects of composition and synthesis routes on the thermoelectric properties of Bi2Te3-based material systems. Material

Synthesis route

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Bi Te Sb Bi2Te3 Bi2Te3 Bi2Te3 Bi2Te3 Bi2Te3 Bi2Te3 Bi2Te3 Bi2Te3 Bi0.5Sb1.5Te3 Bi0.5Sb1.5Te3 Bi0.5Sb1.5Te3 p-type Bi2Te3

– – – Zone melting Zone melting Zone melting SPS SPS SPS Hydrothermal + HP Hydrothermal + HP Nanocrystalline-Ball milling + HP Nanocrystalline-Ball milling + HP Nanocrystalline-Ball milling + HP Melt spinning + SPS

27 27 27 27 87 147 27 87 147 27 227 27 127 227 117

70 500 47 227 234 225 212 222 220 130 148 185 212 201 225

0.00867 x106 2 0.0288  106 940 650 550 1100 800 620 1300 730 1250 750 490 625

42.48 0.5 63.61 48.43 35.59 27.84 49.43 39.42 30 21.97 15.98 42.78 33.70 19.79 31.64

7.87 2.35 24.3 1.81 1.97 2.5 1.09 1.09 1.20 1 0.8 1.12 1.04 1.27 0.8

0.161 0.63  104 0.078 0.80 0.65 0.45 1.35 1.3 1.05 0.65 1.0 1.2 1.4 0.80 1.5

[159] [159] [159] [172] [172] [172] [172] [172] [172] [87] [87] [150] [150] [150] [57]

exfoliated n-type Bi2Te3 2D sheets prepared by SPS, higher than the ingot form (ZT  0.6) [175]. Studies on Bi2Te3 have shown that ZT can be increased by decreasing the grain size and/or increasing the number of pores within the material. However, both can lead to a decrease in the thermal/electrical conductivity; thus, proper control of the microstructure and relative density is needed to overcome this limitation [176]. The addition of metal oxide particles to the bulk matrix via attrition milling can increase the ZT value while decreasing the thermal conductivity [177–179]. Lee et al. synthesized Bi0.5Sb1.5Te3-based nanocomposites by incorporating TiO2 NPs at different vol% via dry milling. They found that the thermal conductivity decreased with an increase in the TiO2 concentration [177]. A ZT of 0.9 was obtained at 67 °C for Bi0.5Sb1.5Te3 nanocomposites with 3 vol% of TiO2 NPs added [177]. Nanocomposites composed of ball-milled (Bi2Te3)0.2(Sb2Te3)0.8 dispersed with amorphous SiO2 NPs of 50-nm size were fabricated by the SPS process. These materials showed a ZT of 1.12 and 1.27 at 30 and 90 °C, respectively, which are higher than those of the undispersed (Bi2Te3)0.2(Sb2Te3)0.8 sample [178]. The power factor of (Bi2Te3)0.2(Sb2Te3)0.8 is increased with a decrease in thermal conductivity due to the scattering of phonons at the phase boundaries and NPs [178]. By mixing 0.4 vol% SiC NPs into the Bi0.3Sb1.7Te3 matrix, the ZT can be increased up to 1.33 at 100 °C [179]. The formation of coherent interfaces between the SiC nanoinclusions and the Bi0.3Sb1.7Te3 matrix increases the Seebeck coefficient while preserving the electrical conductivity, with a consequent enhancement of the mechanical properties. This system also displays a reduced lattice thermal conductivity due to enhanced phonon scattering [179]. Kim et al. studied the effect of dense dislocation on the thermoelectric performance of Bi0.5Sb1.5Te3 [65]. This material exhibited a ZT of 1.86 at 320 K, which could be attributed to the combined effect of grain boundary, point defect, and dislocations generated in an excess of Te-rich melt-spun Bi0.5Sb1.5Te3. The dislocations help decrease the thermal conductivity (0.65 W/m K) with minimal charge-carrier scattering, to obtain a high power factor of 39 lW/cm2 K [65]. Wu et al. studied the effect of doping rare earth elements, such as Ce, Y, and Sm, in R0.2Bi1.8Se0.3Te2.7, achieving a ZT of 1.21 at 140 °C for Y0.2Bi1.8Se0.3Te2.7 [180]. Doping of rare earth elements can reduce the electrical resistivities and provide additional phonon-scattering centers for reducing the thermal conductivities [180]. 6.2.4. AgSbTe2 Ternary polycrystalline AgSbTe2 is found to be effective for low- to medium-temperature thermoelectric applications. Recently, a p-type stoichiometric complex structure of AgSbTe2 was investigated [23–25,58,181]. The phase diagram of the Ag–Sb–Te system is a mixture of AgSbTe2/Sb2Te/Ag2Te metastable phases [181]. AgSbTe2 crystallizes into a rock-salt structure, with the Ag and Sb atoms occupying the Na sites. Techniques such as X-ray diffraction, neutron diffraction, and electron diffraction could not be used to determine the ordering of AgSbTe2, as the Ag, Sb, and Te atoms exhibit similar scattering factors. The domains of AgSbTe2 contain regions of both ordered and disordered Ag and Sb atoms. Because of the complicated chemistry of AgSbTe2, its exact structure, composition, and electronic nature remains to be elucidated [44,58,181]. The effects of various dopants and compositions on the thermoelectric properties of AgSbTe2 material systems are summarized in Table 15. As shown in the table, the power factor increases upon Se doping in the AgSbTe2 material system [182]. The disordered structure of AgSbTe2 plays a key role in the reduction of thermal conductivity [58]. 6.2.5. AgPbmSbTe2+m (LAST) The reaction between AgSbTe2 and PbTe material systems results in the formation of AgPbmSbTe2+m, also known as ‘‘LAST.” AgSbTe2 and PbTe possess a rock-salt structure with p-type characteristics and low lattice thermal conductivities. The complex phase diagram of LAST indicates that its thermoelectric properties are dependent on the processing conditions

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Table 15 Effects of doping and composition on the thermoelectric performance of AgSbTe2. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Sb Te Ag Se AgSbTe2 Ag0.88Sb1.04Te2 Ag0.83Sb1.06Te2 AgSbTe2 Ag0.88Sb1.04Te2 Ag0.83Sb1.06Te2 AgSbSe0.02Te1.98 AgSbSe0.02Te1.98

27 27 27 27 27 27 27 377 377 377 27 327

47 500 1.51 900 264 200 195 242 225 228 270 250

0.0288  106 2 0.63  106 1  106 95 161 200 147 178 172 155 225

63.61 0.5 1.43 8.1  107 6.62 6.44 7.60 8.60 9.0 8.94 11.29 14.06

24.3 2.35 429 2.04 0.5 0.73 0.7 0.38 0.68 0.68 0.58 0.62

0.078 0.63  104 1.00  104 1.19  104 0.4 0.26 0.32 1.57 0.86 0.85 0.58 1.36

[159] [159] [159] [159] [58] [58] [58] [58] [58] [58] [182] [182]

[23,149]. Zhou et al. synthesized LAST nanostructures with enhanced thermoelectric properties [23]. The melting point of LAST is 927 °C. The nanoscale inclusion of minor phases in AgPbmSbTe2+m introduces incoherent or semi-coherent interfaces within the matrix. Metals such as Ag and Pb can create disorder in the rock-salt structure by occupying the Na sites, whereas the chalcogen atoms occupy the Cl sites. The family of AgnPbmMnTem+2n (M = Sb/Bi) exhibit several desirable properties, such as high crystal symmetry, isotropic morphology, low thermal conductivity, and the ability to control the carrier concentration [23–25,44,149]. The effects of various alloying elements on the thermoelectric properties of LAST are summarized in Table 16. The presence of Sb atoms, rather than the Bi atoms, in the AgnPbmMnTem+2n system is found to increase the ZT values [24]. This is attributed to the high thermal conductivity (1.5 W/m K at 327 °C) of Bi alloys [25] compared with Sb alloys (1 W/m K) [19]. These results confirm that not only nanostructuring but also alloying elements and the ordering of atoms are needed to decrease the thermal conductivity of AgnPbmMnTem+2n. A ZT of 2.2 was obtained for LAST in its bulk form [19]. The temperature dependence of the solubility limit of the Ag-doped PbTe/Ag2Te composite [denominated by (AgxPbTe)0.945(Ag2Te)0.055] leads to an increase in carrier concentration with an increase in temperature and a ZT value of 1.4 at 477 °C [79]. The thermoelectric properties of the PbTe/Ag2Te composites are summarized in Table 17. 6.2.6. AgPbmSnnSbTe2+m+n (LASTT) A class of p-type thermoelectric materials, denoted as ‘‘LASTT,” with the general formula AgPbmSnnSbTe2+m+n have been developed with the addition of Sn atoms to LAST. In LASTT, the Sn atoms replace the Pb atoms and the transport properties of LASTT can be tuned by varying the Pb:Sn ratio rather than the concentration of Ag or Sb. The cubic NaCl-type structure of LASTT is stable up to 727 °C. The lattice parameters of LASTT increase with the concentration of Pb atoms in the material. The lattice thermal conductivities of alloys such as AgPb12Sn4Sb0.4Te20, AgPb14Sn6Sb0.4Te24, and AgPb10Sn10Sb0.67Te22 are found to be 40% and 0% higher than that of PbTe at temperatures of 27 and 377 °C, respectively [44,149]. The effects of various alloying elements on the thermoelectric properties of LASTT are summarized in Table 18. Han et al. synthesized LASTT by replacing the Pb atoms entirely with Sn atoms to produce a more eco-friendly material [183]. AgSbTe2 plays an active role in modifying the charge and phonon transport in SnTe, and the ratio of SnTe:AgSbTe2 has a marked effect on the thermoelectric properties of AgSnmSbTem+2. A ZT value of 1 was obtained for AgSn4SbTe5, higher than that for pristine SnTe [183].

Table 16 Effect of alloying elements on the thermoelectric performance of LAST. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Sb Te Ag Pb AgPb10SbTe12 AgPb10SbTe12 AgPb18SbTe20 AgPb18SbTe20 AgPb18SbTe20 Ag0.4Pb22.5SbTe20 Ag0.4Pb22.5SbTe20

27 27 27 27 27 427 27 427 527 27 427

47 500 1.51 1.05 150 290 135 335 370 212 210

0.0288  106 2 0.63  106 0.0481  106 720 135 1850 280 216 119 400

63.61 0.5 1.43 0.0530 16.2 11.35 33.71 31.42 29.57 5.34 17.64

24.3 2.35 429 35.3 1.3 0.8 2.3 1.1 1.08 1.25 0.8

0.078 0.63  104 1.00  104 0.45  104 0.37 0.99 0.43 1.99 2.19 0.12 1.50

[159] [159] [159] [159] [19] [19] [19] [19] [19] [24] [24]

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C. Gayner, K.K. Kar / Progress in Materials Science 83 (2016) 330–382 Table 17 Thermoelectric properties of composite of PbTe/Ag2Te material at 477 °C [79]. Material

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

(Ag0.04PbTe)0.945(Ag2Te)0.055 (Ag1PbTe)0.945(Ag2Te)0.055 (Ag2PbTe)0.945(Ag2Te)0.055) (Ag3PbTe)0.945(Ag2Te)0.055 (Ag4PbTe)0.945(Ag2Te)0.055

205 220 212 220 220

350 250 280 320 280

14.70 12.10 12.58 15.48 13.55

1.25 0.75 0.75 0.75 0.7

0.88 1.21 1.25 1.45 1.4

Table 18 Effect of alloying elements on the thermoelectric performance of LASTT. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Sb Te Ag Pb Sn Ag0.5Pb6Sn2Sb0.2Te10 Ag0.5Pb6Sn2Sb0.2Te10 AgSnmSbTem+2 (m = 2) AgSnmSbTem+2 (m = 2) AgSnmSbTem+2 (m = 4) AgSnmSbTem+2 (m = 4) AgSnmSbTem+2 (m = 5) AgSnmSbTem+2 (m = 5) AgSnmSbTem+2 (m = 10) AgSnmSbTem+2 (m = 10)

27 27 27 27 27 27 427 27 427 27 427 27 427 27 427

47 500 1.51 1.05 1 75 262 75 130 64 135 64 137 50 115

0.0288  106 2 0.63  106 0.0481  106 0.0917  106 1000 210 1800 1100 2400 1200 2400 1200 2800 1300

63.61 0.5 1.43 0.0530 0.091 5.62 14.41 10.12 18.59 9.83 21.87 9.83 22.52 7 17.19

24.3 2.35 429 35.3 66.6 1.4 0.8 1.5 1.45 1.8 1.55 2.18 1.8 3 2.2

0.078 0.63  104 1.00  104 0.45  104 0.41  104 0.12 1.30 0.20 0.89 0.16 0.99 0.13 0.87 0.07 0.54

[159] [159] [159] [159] [159] [182] [182] [183] [183] [183] [183] [183] [183] [183] [183]

Fig. 24. Crystal structure of La3xTe4. Here, the blue and brown spheres represent the La atoms and Te atoms, respectively. Reproduced with permission from [69].

6.2.7. Rare earth chalcogenides Rare earth chalcogenide phases can be observed with the Th3P4-type structure. For instance, Te atoms at the P sites can experience a sixfold coordination with La atoms, producing a distorted octahedral structure, as shown in Fig. 24 [69]. The resulting structure accommodates vacancies on the rare earth site, leading to disorders and distortions in the lattice, and in turn enhanced phonon scattering. This leads to a decrease in the lattice thermal conductivity of 0.4–0.8 W/m K [69]. Snyder et al. studied the effects of carrier and doping concentration on the thermoelectric properties of the La3xTe4 system. They reported that the La3xTe4 system behaves as a degenerate semiconductor with an effective bandgap of 0.9 eV with

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a ZT value of 1.13 at 1000 °C [69,79,127]. Hampl synthesized rare earth-based compositions such as n-type gadolinium selenide (GdSe) and p-type Cu–Ag–chalcogenide (CuAgSe), the thermoelectric properties of which are summarized in Table 19 [184]. 6.2.8. Replacement of Tellurium: PbSe/Bi2S3 Replacing Te with other elements such as Se and S can also be considered as these materials display low thermal conductivities and high Seebeck coefficients. PbSe is inexpensive compared with PbTe, as Se is more abundant than Te [77]. Wang et al. synthesized PbSe doped with Na, yielding a ZT value of >1.2 at 577 °C [77]. Doping elements can generate the same complex band structure in PbSe as observed in PbTe [77]. In their theoretical model, Parker and Singh suggested that a ZT of 2 can be obtained in a heavily doped bulk PbSe if doping can be controlled and the mobility follows the expected trend [132]. A high Seebeck coefficient of 210 lV/K at 602 °C was obtained for a PbSe alloy and a ZT of 1.2 at 602 °C was achieved for K-doped PbSe [78]. The bandgap of K-doped PbSe increases with temperature without any bipolar effects. At elevated temperatures, pinning of the Fermi level is observed due to the heavy band (when the offset value of the two bands is small) [78]. The lattice thermal conductivity of K-doped PbSe decreases from 1.7 at 27 °C to 0.6 W/m K at 602 °C [78]. When Al is doped into the PbSe matrix, the Seebeck coefficient is increased due to an increase in the local DOS near the Fermi level as well as the resonant states [53]. Al doping yields a high ZT, for example, a ZT value of 1.3 at 577 °C for nanosized Al-doped PbSe [53]. Other dopants such as B, Ga, In, and Tl can also enhance the thermoelectric properties of PbSe. Of these, Tl dopants can modify the band structure of the resulting material to yield a high Seebeck coefficient [76]. Nanostructured Na-doped PbSe embedded with endotaxial inclusions of 1–4% MSe (M = Ca/Sr/Ba) is synthesized via SPS. Due to the presence of mesoscale grains and possibly the hierarchical structuring of p-type PbSe-MSe systems, the lattice thermal conductivity decreases significantly [185]. ZT values of 1.2, 1.3, and 1.3 are obtained at 650 °C for Pb0.95Ca0.04Na0.01Se, Pb0.97Sr0.02Na0.01Se, and Pb0.96Ba0.03Na0.01Se, respectively [185]. Bi2S3 contains a highly anisotropic one-dimensional orthorhombic structure. Bi atoms form a highly distorted octahedron, and three short covalent BiAS bonds form Bi2S3 chains. Its unique structure includes infinite chains directed along the crystallographic c-axis with a needle-like morphology [186]. Bi2S3 exhibits a high Seebeck coefficient and low thermal conductivity at room temperature, and it is only limited by its low electrical conductivity. The addition of BiCl3 to Bi2S3 can increase the carrier concentration and electrical conductivity without altering the thermal conductivity [186]. For example, the carrier concentration is found to increase from 3.7  1016 to 2.6  1019 carriers/cm3 upon the addition of 0.5 mol% BiCl3 to Bi2S3 [186]. The thermoelectric properties of Bi2S3 and its components are listed in Table 20. 6.2.9. Copper ion liquid-like thermoelectric Cu2Se The concept of PGEC is found to be highly effective in the case of skutterudites, clathrates, and other intermetallics [12,89,145,163,164]. Liu et al. proposed a Cu2Se system with a low thermal conductivity of 0.4–0.5 W/m K at 727 °C [88]. Subsequently, Cu2xSe was developed with liquid-like behavior of copper ions around a crystalline sublattice of Se [88]. A liquid is similar to a glass due to the lack of long-range periodicity in both. Thermal transport through a disordered

Table 19 Thermoelectric properties of rare earth metal-based materials [184]. Seebeck coefficient (lV/K)

Operating temperature (°C)

300 400 500 600 700 800

Electrical conductivity (S/cm)

GdSe

CuAgSe

GdSe

CuAgSe

221 249 273 292 302 310

191 221 249 270 285 291

163 136 119 106 98 93

227 172 136 112 97 87

Table 20 Thermoelectric properties of Bi2S3 and its components. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Bi S Se Bi2S3 Bi2S3 Bi2S3 with 0.5 mol% BiCl3 Bi2S3 with 0.5 mol% BiCl3

27 27 27 27 377 27 487

70 – 900 350 325 105 233

0.00867  106 5  1018 1.0  106 10 13 615 107

42.48 – 0.81  106 1.22 1.37 6.78 5.80

7.87 0.296 2.04 1.3 0.85 1.38 0.76

0.161 – 0.0119  106 0.028 0.1 0.14 0.6

[159] [159] [159] [186] [186] [186] [186]

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Fig. 25. Crystal structure of Cu2Se with proposed possible Cu distribution at (1/4, 1/4, 1/4), (1/3.17, 1/3.17, 1/3.17), and (1/2, 1/2, 1/2). Reproduced with permission from [54].

structure can be viewed as the random diffusion of energy due to atomic jumps from one rest position to another, and the time lag for this process is longer in solids (crystal/glass) than in liquids. Thus, a disorder in the structure can disrupt the propagation of phonons, and a liquid inhibits the propagation of transverse waves [88]. The crystal structure of Cu2Se is depicted in Fig. 25. The unique combination of monoatomically ordered Se layers and disordered Cu layers in the crystal structure of b-Cu2Se polycrystals is key to reducing the lattice thermal conductivity (0.4–0.5 W/m K); a ZT of 1.6 at 700 °C is noted for this system [54]. The thermoelectric properties of Cu2Se are summarized in Table 21. The power factor of Cu2Se can be increased by increasing its electrical conductivity. Riha et al. synthesized Cu2Se NPs via the hot injection method. They found that, under ambient conditions, the electronic nature of the as-synthesized Cu2Se NPs can be varied from semiconducting to metallic [187]. 6.2.10. Si–Ge alloys Si–Ge alloys are intermetallic materials with good mechanical properties, with potential applications in thermoelectric devices. Si–Ge alloys possess a diamond cubic structure with two interpenetrating FCC lattices separated by an a/4 unit along each axis of the cell where the lattice parameter is varied as a function of composition. The melting point of Si–Ge alloy is 1027 °C, indicating their operation at high temperatures [71,188]. The major drawback of Si–Ge alloys is their high thermal conductivities at room temperature. Wang et al. synthesized fine-grained Si–Ge polycrystals with low thermal conductivities in the order of 2.5–3 W/m K, attributed to the boundary scattering of phonons [160]. The thermoelectric properties of Si–Ge alloys are summarized in Table 22. In a Si–Ge alloy, the addition of Ge to the Si matrix increases the ZT value due to the reduction in thermal conductivity. This reduction in thermal conductivity can be attributed to an enhanced scattering of phonons from the Ge atoms, due to the random distribution of Si and Ge atoms within the alloy [188]. The methodology and processing parameters are known to affect the thermoelectric properties (particularly, the thermal conductivities) of alloys [72,188,189]. A large difference is noted between the mean free paths of electrons and phonons in a highly doped nanostructured Si-Ge alloy: the electron mean free path was 5 and that of phonons varied from 2 to 300 nm [189]. Zhu et al. synthesized a Si95Ge5 system and found that a nanosized grain in Si95Ge5 causes a reduction in thermal conductivity by a factor of 2 compared with the bulk sample. This reduction in thermal conductivity is achieved by limiting the mean free path of phonons with wavelengths of >1 nm. The thermal conductivity of the Si–Ge alloy is found to decrease with increasing molar fraction of Ge [139]. Wang et al. developed an n-type nanostructured Si–Ge bulk alloy yielding a ZT value of 1.3 at 900 °C. The increased ZT of n-type Table 21 Thermoelectric properties of Cu2Se and its components. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Cu Se b-Cu2Se b-Cu2Se b-Cu1.98Se b-Cu1.98Se Cu2Se1.01 (Mechanical alloying + HP) Cu2Se1.01 (Mechanical alloying + HP)

27 27 127 727 127 727 27 700

6.5 900 100 300 60 190 75 250

0.596  106 1.0  106 1000 129 1666 333 1428 181

25.18 0.81  106 10 11.61 6 12 8.03 11.31

401 2.04 1.1 0.74 1.5 1.1 0.97 0.7

0.0018 0.0119  106 0.36 1.5 0.16 1.1 0.24 1.6

[159] [159] [88] [88] [88] [88] [54] [54]

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Table 22 Thermoelectric properties of n- and p-type Si–Ge alloys and their constituents. Material

Semiconducting nature

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Si Ge Nanostructured Nanostructured Nanostructured Bulk Si0.78Ge0.2 Bulk Si0.78Ge0.2 Bulk Si0.78Ge0.2 Nanostructured Nanostructured Nanostructured

– – n-type n-type n-type n-type n-type n-type p-type p-type p-type

27 27 27 600 900 27 527 927 27 600 900

148 280 110 264 242 108 233 242 110 185 270

111 357 750 400 520 1176 537 561 880 520 280

2.43 27.98 9.07 27.87 30.45 13.71 29.15 32.85 10.64 17.80 20.41

148 59.9 2.5 2.5 2.8 4.45 4.0 4.2 2.3 2.4 2.4

0.00049 0.014 0.11 0.97 1.3 0.092 0.58 0.93 0.14 0.65 0.95

[159] [159] [160] [160] [160] [188] [188] [188] [189] [189] [189]

SiGe SiGe SiGe

Si80Ge20 Si80Ge20 Si80Ge20

nanostructured Si–Ge bulk alloys can be ascribed to the significant reduction in their thermal conductivity, caused by a high rate of phonon scattering along with the increased density of nanograin boundaries [160]. The ZT of n-type nanostructured Si–Ge bulk alloys (1.3) increases by 40% compared with that of pristine Si–Ge alloys (0.7) [188]. Porous nanograined materials exhibit low ZTs compared with their bulk counterparts, which is due to the comparatively lower electrical conductivities of the former. Lee et al. proposed a model to elucidate the effect of nanoscale porosity on the transport properties of electrons and phonons [72]. According to this model, the electrons scattered from the pores can be considered as quantum systems, whereas phonon transport follows a classical path. Thus, the charge carriers are scattered to a greater extent in nanograin materials than in their bulk forms as a large number of scattering sites are available. In porous nanograined materials, low thermal conductivities are observed due to the enhanced phonon scattering, as well as high Seebeck coefficients primarily caused by energy filtering [72]. For p-type nanostructured bulk Si–Ge alloys, a ZT of 0.95 is obtained, which is 90% higher than that for Si–Ge alloys used in space flight missions and 50% higher than that for p-type Si–Ge alloys [189]. Bathula et al. synthesized an n-type Si80Ge20 nanostructured alloy via high-energy ball milling followed by SPS, yielding a ZT value of 1.5 at 900 °C [59]. High-energy ball milling for 90 h and the formation of an additional phase of SiP nanoprecipitates during the subsequent SPS process lead to a decrease in the thermal conductivity of 2.3 W/m K at 900 °C. The SiP nanoprecipitates create a compositional difference between the grain and the grain-boundary region, ultimately enhancing the scattering of phonons and in turn the ZT value. 6.2.11. HH alloys HH alloys constitute the intermetallic class of compounds, with potential applications in high-temperature thermoelectric devices. HH alloys contain a MgAgAs-type cubic structure with three interpenetrating FCC sublattices, each occupied by X, Y, and Z atoms. Here, X denotes a transition metal, a noble metal, or a rare earth element (e.g., Ti/Hf/Zr); Y a transition metal or noble metal (e.g., Co/Mn); and Z a metalloid or a metal (e.g., Sb). This cubic structure contains 18 valence electrons and a complex band structure. The covalent bonding between two elements chiefly determines determining the electronic properties; further, bonding configurations have been implicated in phase stability and the occurrence of a bandgap [190]. The presence of vacant atomic sites in an HH alloy can lead to the formation of narrow bands, resulting in a d-orbital hybridization and semiconducting character [191]. Most of the HH alloys exhibit narrow bandgaps ranging between 0.1 and 0.3 eV [190,191]. HH alloys exhibit high Seebeck coefficients (up to 300 lV/K) and high electrical conductivity (103–104 S/cm) at room temperature, but they are significantly limited by their relatively high thermal conductivity (10 W/m K). These materials also display high melting point (1100–1300 °C) and good thermal stability (up to 1000 °C) [44,51,190]. The atomic disorder at the transition metal sites can reduce the lattice thermal conductivity due to induced mass fluctuations and strain field effects in HH alloys [147,148]. The effects of composition on the thermoelectric properties of HH alloys are listed in Table 23. Sb-doped TiNiSn alloys exhibit a high power factor of 70 lW/cm K2 at 377 °C. However, the drawback of these alloys is their high thermal conductivity of 10 W/m K, with a low ZT of 0.45 at 377 °C [193]. Shutoh and Sakurada synthesized a Ti0.5(Zr0.5Hf0.5)0.5NiSn1YSbY alloy, with a ZT of 1.5 at 527 °C [51]. The substitution of Ti with Zr and Hf reduced the thermal conductivity to as low as 3.1 W/m K at 527 °C, thus leading to an increased Seebeck coefficient [51,81]. Yu et al. developed Hf0.6Zr0.4NiSn0.98Sb0.02 without Ti alloying by levitation melting followed by SPS, yielding a ZT of 1.0 at 727 °C [86]. Chen et al. prepared Hf0.75Zr0.25NiSn0.99Sb0.01 with a ZT of 1 at 600–700 °C [194]. Zr0:25Hf0:75NiSn alloys prepared by highenergy ball milling and SPS exhibit increased ZT values compared to their bulk forms synthesized by arc melting [192]. The key advantages of Zr0:25Hf0:75NiSn alloys are the simultaneous decrease in thermal conductivity and increase in the Seebeck coefficient [192]. Recent studies have shown that the thermal conductivity may decrease due to mass fluctuation and strain field fluctuation. The irregular shapes of micro- and nanostructures in an arc-melted solid solution of Ti0.37Zr0.37Hf0.26NiSn led to ZT values of 1 at 452 °C and 1.5 at 352 °C, and a low thermal conductivity of 2.3 W/m K. This can be attributed to the scattering of phonons by mass and strain fluctuation effects [195]. The lattice thermal conductivity

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C. Gayner, K.K. Kar / Progress in Materials Science 83 (2016) 330–382 Table 23 Thermoelectric performance of HH alloys and their components. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Ti Hf Ni Sn Sb Zr Ti0.5(Zr0.5Hf0.5)0.5NiSn1YSbY (Y  0) Ti0.5(Zr0.5Hf0.5)0.5NiSn1YSbY (Y  0.002) Ti0.5(Zr0.5Hf0.5)0.5NiSn1YSbY (Y  0.006) Ti0.5(Zr0.5Hf0.5)0.5NiSn1YSbY (Y  0) Ti0.5(Zr0.5Hf0.5)0.5NiSn1YSbY (Y  0.002) Ti0.5(Zr0.5Hf0.5)0.5NiSn1YSbY (Y  0.006) Hf1xZrxNiSn0.98Sb0.02 (x  0.40) by HP Hf1xZrxNiSn0.98Sb0.02 (x  0.40) by SPS Hf1xZrxNiSn0.98Sb0.02 (x  0.40) by HP Hf1xZrxNiSn0.98Sb0.02 (x  0.40) by SPS Hf1xZrxNiSn0.98Sb0.02 (x  0.50) by HP Hf1xZrxNiSn0.98Sb0.02 (x  0.50) by SPS Hf1xZrxNiSn0.98Sb0.02 (x  0.50) by HP Hf1xZrxNiSn0.98Sb0.02 (x  0.50) by SPS Zr0.5Hf0.5Ni0.8Sn0.99Sb0.01 Zr0.5Hf0.5Ni0.8Sn0.99Sb0.01 + ZrO2 NPs (3 vol%) Zr0.5Hf0.5Ni0.8Sn0.99Sb0.01 + ZrO2 NPs (6 vol%) Zr0.5Hf0.5Ni0.8Sn0.99Sb0.01 + ZrO2 NPs (9 vol%) Zr0.5Hf0.5Ni0.8Sn0.99Sb0.01 + ZrO2 NPs (0 vol%) Zr0.5Hf0.5Ni0.8Sn0.99Sb0.01 + ZrO2 NPs (3 vol%) Zr0.5Hf0.5Ni0.8Sn0.99Sb0.01 + ZrO2 NPs (6 vol%) Zr0.5Hf0.5Ni0.8Sn0.99Sb0.01 + ZrO2 NPs (9 vol%) Nanostructured Zr0.25Hf0.75NiSn Nanostructured Zr0.25Hf0.75NiSn

27 27 27 27 27 27 27 27 27 527 527 527 27 27 527 527 27 27 527 527 27 27 27 27 627 627 627 627 27 500

9.1 5.5 19.5 1 47 8.9 309 250 210 300 280 243 98 92 148 165 78 92 125 154 87 90 96 96 145 137 145 160 130 187

0.0234  106 0.0312  106 0.143  106 0.0917  106 0.0288  106 0.0236  106 192 400 909 500 666 714 2000 2350 1250 1600 1700 2250 900 1500 1739 1724 1626 1515 1052 970 1020 854 1910 1110

1.93 0.943 54.37 0.0917 63.61 1.86 18.33 25 40.08 45 52.21 42.16 19.20 19.89 27.38 43.56 10.34 19.04 14.06 35.57 13.16 13.96 14.98 13.96 22.11 18.20 21.44 21.86 33 38

21.9 23 90.7 66.6 24.3 22.7 3.4 3.25 3.3 3.1 2.9 2.8 4.6 5 3.9 4.1 3.8 4.65 3.3 3.9 4.5 4.25 3.9 3.4 3.7 3.6 3.4 2.7 3.8 2.6

2.56  103 1.23  103 0.0179 4.13  105 0.078 2.47  103 0.16 0.23 0.36 1.16 1.5 1.20 0.12 0.11 0.56 0.85 0.081 0.12 0.34 0.72 0.087 0.098 0.11 0.12 0.53 0.45 0.56 0.72 0.27 1.1

[159] [159] [159] [159] [159] [159] [51] [51] [51] [51] [51] [51] [86] [86] [86] [86] [86] [86] [86] [86] [73] [73] [73] [73] [73] [73] [73] [73] [192] [192]

rapidly increased with increasing Nb content in a FeVSb alloy. At room temperature, the observed values of jlattice for FeVSb and FeV0.6Nb0.4Sb alloys were 12.2 and 5.6 W/m K, respectively [196]. Makongo et al. stated that the synthesis routes, processing conditions, and microstructures are key parameters determining the sensitivity of electrical conductivity, thermal conductivity, and carrier mobility [197]. In a mechanically alloyed sample, a decrease in grain size leads to a significant decrease in carrier mobility and lattice thermal conductivity as well. In an SPS-processed sample, grain growth results in a low lattice thermal conductivity and moderate carrier mobility [197]. Nanostructuring via ball milling and HP is proven to be effective in increasing the ZT up to 1 at 600–700 °C for n-type Hf0.75Zr0.25NiSn0.99Sb0.01 HH alloys; furthermore, grain sizes <100 nm are proven effective in further increasing the ZT value [198]. Yan et al. reported a 60% increase in ZT for p-type HH nanopowder prepared by ball milling and HP [199]. Upon substituting excess Ni, TiNi1+xSn (x = 0–0.15) prepared by microwave and then SPS processes showed decreased thermal conductivity and increased power factor [200]. Chen et al. prepared a nanocomposite by dispersing nanophase inclusions of ZrO2 into the Zr0.5Hf0.5Ni0.8Pd0.2Sn0.99Sb0.01 matrix by SPS [73]. The thermal conductivity of this nanocomposite is decreased due to the scattering of phonons by the ZrO2 NP aggregates at the grain boundaries, which act as scattering centers for phonons [73]. However, the Seebeck coefficient and electrical resistivity are found to increase, due to the potential barrier scattering. Upon grain refinement and embedding of NPs, n-type Hf0.6Zr0.4NiSn0.995Sb0.005 HH alloys and p-type Hf0.3Zr0.7CoSn0.3Sb0.7/nano-ZrO2 composites exhibited increased ZT values of 1.05 and 0.8 at 627 and 727 °C, respectively [201]. According to the ASSET approach, the thermopower and electrical conductivity can be simultaneously increased in a bulk HH alloy embedded with Hf inclusions [48]. The approach was successful in the case of HH bulk nanocomposites ((1x)ZrzHf1zNiSn1ySby + xZrzHf1zNi2Sn1yBiy, 0 6 z, y 6 1, 0 6 x 6 0.1), which increased the power factor from 10 to 47 lW/K2 from 454 to 527 °C [48]. Sahoo et al. synthesized the bulk HH matrix of Ti0.5Hf0.5CoSb0.9Sn0.1 embedded with different mole fractions of Hf via solid-state reactions. They reported the efficacy of this method in increasing the thermopower and minimizing the decrease in electrical conductivity [157]. A heavily doped HH alloy is subject to a hole-culling effect that tunes its mobility without affecting its thermal conductivity. The reduction in effective carrier density can increase the Seebeck coefficient [157]. For HH alloys, multiatom substitution is found to be more effective in enhancing their thermoelectric properties. The highest power factors at 427 °C are reported for multiatom-filled HH alloys, as shown in Table 24. Multiatom filling and high-temperature annealing are known to reduce the thermal conductivity significantly and enhance the electrical properties, yielding high ZT values

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Table 24 Thermoelectric properties of HH alloys [52]. Material

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT (at 427 °C)

(Er0.05(Ti0.38Zr0.32Hf0.30)0.95)33Ni34 (Sn0.997Sb0.003)33 (Ti0.32Zr0.30Hf0.30Ta0.08)34Ni34Sn32 (Ti0.35Zr0.35Hf0.30)33(Ni0.998Cu0.002)35(Sn0.988Sb0.012)32 (Ti0.30Zr0.35Hf0.35)29Ni33(Sn0.994 Sb0.006)38 (Gd0.32(Ti0.38Zr0.32Hf0.30)0.68)33 Ni34(Sn0.997Sb0.003)33 (Zr0.50Hf0.50)33Ni34(Sn0.997Sb0.003)33

262 260 274 275 275 271

800 819 800 800 813 826

54.91 55.36 60.06 60.50 61.43 60.66

2.7 2.8 2.8 2.8 2.9 3.1

1.42 1.38 1.50 1.51 1.51 1.37

[52]. In the near future, approaches such as nanostructuring and bandgap engineering can yield high ZT values >2 by optimizing the thermoelectric properties of the HH alloys. 6.2.12. Clathrates Clathrates possess a cage-like crystal structure, formed from the guest atoms in the crystal lattice, exhibiting low thermal conductivity. The families of clathrates tetrahedrally coordinate with Al, Si, Ga, Ge, or Sn forming an open framework. Based on this coordination, the system is subdivided into type I (A8E46), type II (A24E136), and type III [91,92,202]. The structure of A8E46 is depicted in Fig. 26. It is composed of covalently bound Ga–Ge with the cages containing rattling cations [92]. Cations are attached to the anion frameworks of tetrakaidecahedral and dodecahedral polyhedra. In these polyatomic compounds, one component forms a cage structure around another [92,202]. Clathrates have complex structures with more than one phase present, although the semiconducting phase alone is desirable for thermoelectric applications. Clathrates behave as ‘‘glass-like” materials with electronic properties than can be varied with doping [92,202,203]. The guest atoms function as a ‘‘rattler” inside the host matrix, scattering lattice phonons and in turn reducing the thermal conductivity [204]. However, the concept of ‘‘rattling” is debatable due to the presence of lattice disorder and point defects, which are known to reduce the lattice thermal conductivity. The effects of composition on the thermoelectric properties of clathrates are summarized in Table 25. The Ba8Ga16Ge30 crystal synthesized by the Czochralski method displayed a ZT of 1.35 at 627 °C, which can be extrapolated to reach a ZT of 1.63 at 827 °C [55]. For type III clathrates, such as Ba8Ga15Ge85, a ZT of 1.25 at 627 °C and a reduced thermal conductivity of 0.8 W/m K were reported [205]. Tang et al. synthesized YbxBa8xGa16Ge30 via a combination of melting and SPS, with a ZT value of 1.1 [203]. The Yb atoms added fill the voids of Ga and Ge in YbxBa8xGa16Ge30; further, the increase in Yb content in the matrix increases the electrical conductivity but decreases the Seebeck coefficient and lattice thermal conductivity [203]. Zhang et al. found that synthesis methods such as SPS can have a significant effect on the thermoelectric performance in a single crystal of Ba8Au5.3Ge40.7 prepared by the Bridgman method [91]. This crystal exhibits a p-type behavior and a low thermal conductivity with a high Seebeck coefficient. The ZT of the single-crystal Ba8Au5.3Ge40.7 system increases from 0.3 at 227 °C to 0.9 at 407 °C upon SPS processing [91]. Toberer et al. reported an extremely low lattice thermal conductivity of 0.14 W/m K at 727 °C in Ba8Ga16Ge30 with a relatively low ZT value of 0.8 [92]. For type I Ba8Ga16Ge30 formed by the cross-substitution of the framework elements, a power factor of 15 lW/cm K2 was obtained, ascribed to the change in the carrier scattering mechanisms. A ZT of 1.2 at 727 °C was reported for a polycrystalline Ba8Ni0.31Zn0.52Ga13.06Ge32.2 material system [206]. The framework element substitution introduces ionized impurities and

Fig. 26. Type I clathrate, Ga–Ge covalently bonded structure with cages containing cations. The anion framework of tetrakaidecahedral and dodecahedral polyhedra are indicated in blue with red cations. Reproduced with permission from [92].

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C. Gayner, K.K. Kar / Progress in Materials Science 83 (2016) 330–382 Table 25 Thermoelectric properties of clathrates. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Ba Ge Ga Ba8Ga16Ge30 (Czochralski method) Ba8Ga16Ge30 Ba8Ga16Ge30 (Melt & SPS) Yb0.3Ba7.7Ga16Ge30 Yb0.5Ba7.5Ga16Ge30 Yb0.7Ba7.3Ga16Ge30 Yb1.0Ba7Ga16Ge30 Ba8Ga16Ge30 Yb0.3Ba7.7Ga16Ge30 Yb0.5Ba7.5Ga16Ge30 Yb0.7Ba7.3Ga16Ge30 Yb1.0Ba7Ga16Ge30 Ba8Au5.3Ge40.7 (Single crystal – Bridgman method) Ba8Ga15Ge85 (Arc melting) Ba8Ga15Ge85

27 27 27 27 627 27 27 27 27 27 677 677 677 677 677 227 27 627

12.5 280 – 45 180 69 54 60 50 41 218 205 194 170 140 340 57 148

0.03  106 357 0.0678  106 1530 570 520 492 591 943 740 200 280 330 520 380 40 833 550

4.68 27.98 – 3.09 18.46 2.47 1.43 2.12 2.35 1.24 9.50 11.76 12.41 15.02 7.44 4.60 2.70 12.04

18.4 59.9 40.6 1.8 1.25 2.04 1.55 1.5 1.68 1.9 1.54 1.35 1.16 1.5 1.65 0.7 0.8 0.8

7.64  103 0.014 – 0.051 1.35 0.036 0.027 0.042 0.041 0.019 0.58 0.82 1.1 0.95 0.42 0.3 0.024 1.25

[159] [159] [159] [55] [55] [203] [203] [203] [203] [203] [203] [203] [203] [203] [203] [91] [205] [205]

lattice defects into the material, thus inducing the scattering of phonons [206]. For Ba8Ni0.22Zn7.22Ge37.12Sn1.44 prepared by Ni substitution, a ZT value of 0.87 at 557 °C was reported [207]. Sr8Ga16-xGe30-y synthesized by a solid-state reaction followed by SPS exhibited a thermal conductivity of <1 W/m K at 27 °C with an increased power factor of 12 lW/cm K2; further, the power factor increased with a decrease in the Ga/Ge ratio [208].

6.2.13. Other emerging systems Apart from those discussed previously, a few systems with comparable thermoelectric properties have emerged. Examples include b-Zn4Sb3, Yb14MnSb11, Ti9BiTe6, Tl2SnTe5, In4Se3d, and NaCo2O4. Among these materials, b-Zn4Sb3 has gained interest as it can operate at moderate temperatures and exhibits comparatively low thermal conductivity due to the PEGC behavior [146]. The b-Zn4Sb3 structure is composed of one Zn site and two independent Sb sites. The material decomposes to ZnSb and Zn before reaching its melting point of 568 °C. It exhibits a p-type semiconducting behavior, with the Zn atom playing the dual role of thermopower enhancer and electron donor. A ZT of 1.3 at 400 °C was reported for b-Zn4Sb3, and doping was found to have no effect on the ZTs [146]. The thermoelectric properties of the emerging material systems are summarized in Table 26. The other class of anisotropic material, Tl9BiTe6, comprises ternary compounds with a semimetallic behavior. The two main advantages of Tl9BiTe6 are the low thermal conductivity due to the large average mass of its constituent atoms in a unit cell and the low melting point [140]. The crystal structure of Tl9BiTe6 exhibits disorder at the heavy atom sites, which are occupied by Bi and Tl. Its low lattice thermal conductivity (0.4–0.5 W/m K) and Seebeck coefficient of >350 lV/K yield a

Table 26 Thermoelectric properties of emerging material systems. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

Sb Bi Te Ti Se Na Yb b-Zn4Sb3 Yb14MnSb11 Ti9BiTe6 Tl2SnTe5 In4Se3d (d = 0.65) NaCo2O4 SnSe

27 27 27 27 27 27 27 400 927 227 127 432 727 650

47 70 500 9.1 900 2 30 220 180 390 200 280 150 350

0.0288  106 0.00867  106 2 0.0234  106 1.0  106 0.21  106 0.0351  106 280 185 58 280 200 550 100

63.61 42.48 0.5 1.93 0.81  106 0.84 31.59 13 6 9.0 10 15.6 12.5 10

24.3 7.87 2.35 21.9 2.04 141 34.9 0.7 0.8 0.4 0.5 0.74 1.5 0.35

0.078 0.161 0.63  104 2.56  103 0.0119  106 0.17  105 0.0271 1.3 1.0 1.2 0.8 1.48 0.8 2.62

[159] [159] [159] [159] [159] [159] [159] [146] [209] [140] [210] [60] [211] [64]

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Table 27 Thermoelectric properties of AgxCuyTlTe1+z at 27 °C [212]. Material

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm.K2)

Thermal conductivity (W/m.K)

ZT

Ag0.62Cu0.30TlTe1.05 Ag0.68Cu0.26TlTe1.05 Ag0.64Cu0.28TlTe1.05

350 370 320

140 50 80

17.15 6.84 8.19

0.3 0.3 0.3

1.7 0.68 0.81

Table 28 Thermoelectric properties of some of the Zintl materials [61]. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Yb14MnSb11 Yb14MnSb11 Yb14Mn0.67Zn0.33Sb11 Yb14Mn0.67Zn0.33Sb11 Yb14Mn0.33Zn0.67Sb11 Yb14Mn0.33Zn0.67Sb11 Yb14Mn1Al0.00Sb11 Yb14Mn0.8Al0.2Sb11 Yb14Mn0.4Al0.6Sb11

27 927 27 927 27 927 927 927 927

43 190 48 240 45 190 200 230 230

500 185 555 212 833 222 212 208 131

0.92 6.67 1.27 12.21 1.68 8.0 8.48 11 6.92

0.83 0.72 0.76 0.83 0.83 0.65 0.7 0.74 0.62

0.03 1.10 0.050 1.7 0.060 1.47 1.48 1.76 1.33

ZT value of 1.2 at 227 °C [140]. Tl2SnTe5 is tetragonal in structure with columns of Tl ions along the crystallographic c-axis. The large interatomic distance makes Tl2SnTe5 a promising candidate for thermoelectric application [210]. It exhibits a low thermal conductivity and a ZT of 0.8 at 127 °C. Rhyee et al. synthesized a binary crystalline n-type In4Se3d with a ZT of 1.48 is obtained at 432 °C, arising from the high Seebeck coefficient of >350 lV/K and low thermal conductivity of 0.74 W/m K [60]. Brun et al. proposed a new compound with the general formula ‘‘AgxCuyTlTe1+z,” the thermoelectric properties of which are presented in Table 27 [212]. At room temperature, it exhibits a ZT value of 1.7 with a markedly low thermal conductivity of 0.3 W/m K [212]. The ultralow thermal conductivity and enhanced electrical transport lead to a ZT of 1 and 1.15 at 407 °C for 4 mol% Pbdoped and 2 mol% Bi-doped AgSbSe2, respectively [213]. These values are found to be 150% and 190% higher than that of pristine samples. (AgCrSe2)0.75(CuCrSe2)0.25 exhibits significantly increased ZT of 1.4 at 427 °C, which can be attributed to its low thermal conductivity of 1 W/m K [214]. Zhao et al. synthesized a SnSe cubic system with a ZT of 2.62 at 923 K, primarily resulting from its ultralow lattice thermal conductivity of 0.35 W/m K [64]. A change in the distorted structure of SnSe can affect the ZT values, which vary from 2.3 to 0.8. The layered SnSe possesses a distorted rock-salt structure, which not only reduces the lattice thermal conductivity but also maintains a high Seebeck coefficient of 350 lV/K [64]. The high-temperature intermetallics compound Yb14MnSb11 contains a Zintl phase of A14MPn11-type (A = alkaline earth or rare earth metal, M = transition or main group metal, and Pn = pnicogen) structure [209]. The low thermal conductivity of Yb14MnSb11 can be attributed to the complexity in the structures and a large atomic mass, restricting phonons by reducing the fraction of atomic vibrational modes that drop to carry heat efficiently [146,209]. A ZT value of 1 at 927 °C was reported for this material, which is expected to replace p-type Si–Ge alloys in the near future due to the superior thermoelectric properties of the former [209]. Due to the complex cubic structure of Yb14MnSb11, its electronic behavior such as semiconducting, metallic, or semimetallic nature cannot be specified as a function of temperature [209,215]. Synder et al. developed a Zintl material, A14MPn11, which acts as a small bandgap semiconductor [61]. A14MPn11 contains cationic sites that can tune the carrier concentration and in turn induce disordered scattering [61]. The thermoelectric performance of some of the Zintl materials is summarized in Table 28. A high ZT of 1.76 at 927 °C was obtained for Yb14Mn0.8Al0.2Sb11, resulting from the low lattice thermal conductivity [61]. Metal oxides are known for their high thermal and chemical stabilities, with potential application in thermoelectric devices. They exhibit low thermal conductivities, which can eventually lead to high ZT [216]. In general, most metal oxides are poor conductors of electricity due to their low charge-carrier mobility. However, NaCo2O4, which has a layered structure, Table 29 ZTs of layered Co-oxides at different operating temperatures [211]. Operating temperature (°C)

27 300 500 700

ZT NaxCoO2 crystal

Bi2Sr2Co2Oy crystal

Ca3Co4O9 crystal

NaxCoO2 ceramic

NaxCoO2 ceramic

0.05 0.25 1.0 –

0.07 0.17 0.4 1.14

0.07 0.2 0.45 0.88

0.07 0.28 0.57 0.8

0.17 0.4 0.6 0.8

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is excluded from this group due to its metal-like conductivity. In the structure of NaCo2O4, the CoO2 layer is formed by the edge-sharing distorted octahedral, and the Na+ ions randomly occupy 50% of the interlayer sites [211,216]. The sheets of CoO2 act as electron-transport layers and the intercalated Na+ ions as the phonon-scattering region, ultimately yielding a low thermal conductivity of 3 W/m K at room temperature and 1.5 W/m K at 727 °C [211,216]. A power factor of 50 lW/cm K2 is obtained for the single-crystal NaCo2O4 at 27 °C [216]. A ZT of 0.8 at 727 °C is obtained for polycrystalline NaCo2O4 [211]. Promising alternative thermoelectric materials include p-type metal oxide materials, such as Bi-doped Ca3Co4O9, (Bi, Pb)2Sr2Co2O8, TlSr2Co2Oy, and (Hg, Pb)Sr2Co2Oy. The ZTs obtained for various p-type metal oxides are listed in Table 29. Sui et al. fabricated a textured Bi0.875Ba0.125CuSeO via hot forging [62], which was found to increase the carrier mobility. This in turn increased the electrical conductivity and the power factor from 6.3 lW/cm K2 (before hot forging) to 8.1 lW/cm K2 at 650 °C (after hot forging). The highest ZT obtained for oxygen-containing thermoelectric materials can be increased from 1.1 to 1.4 with the use of texturing processes [62]. The major drawbacks of metal oxide materials are their high contact resistances at the oxide/metal interfaces and the cracking/exfoliation arising from the difference in their thermal expansion coefficients [217]. Rhyee et al. invented a new class of materials derived from CeTe3, with the formula (R1a, Ra0 ) (T1b, Tb0 )3y [218]. The thermal conductivity of these materials range from 2 to 2.5 W/m K within a temperature range of 27–427 °C. A power factor of 22 lW/cm K and a ZT of 0.55 are reported at room temperature; in addition, the thermoelectric properties of these materials are found to vary with their composition [218]. Rhyee et al. proposed a new class of materials with the stoichiometric formula AaRbG3±n [219]. These materials exhibited high electrical conductivities in the range of 1000–4000 S/cm with low lattice thermal conductivities at 27–127 °C [219]. Sharp proposed a p-type semiconducting material, X2YZ5, with a ZT of 0.65 at 27 °C, which is higher than that obtained for Bi2Te3 crystals [220]. The thermoelectric properties of this material can be enhanced via annealing, which reduced the concentration of p-type defects. Tl2SnTe5 exhibits a Seebeck coefficient of 165 lV/K at 27 °C with an electrical conductivity of 250 S/cm and a thermal conductivity of 0.7 W/m K, which is less than that reported for Bi2Te3 [220]. A high ZT of 1.4 at 527 °C is obtained for Bi-doped Mg2.16(Si0.4Sn0.6)1yBiy solid solutions prepared by SPS [63]. Doping of Bi increases the density of electrons, in turn increasing the electrical conductivity by up to 1–2 orders of magnitude. The highest power factor obtained for a solid solution with a Bi concentration of 0.03 wt% is 42 lW/ cm K2 [63]. Mg1Ag1Sb1 exhibits ZT values of 0.3 and 0.7 at 27 and 210 °C, respectively, with a power factor of 16 lW/cm K2 at 67 °C (this value is 40% of the Bi2Te3). Mg1Ag1Sb1 displays a thermal conductivity of 1.14 W/m K at 67 °C, which is 80% that of p-type Bi2Te3-Sb2Te3 and 50–60% that of pristine Bi2Te3 [221]. The thermoelectric properties of Mg1 Ag1 Sb1 at different operating temperatures are summarized in Table 30. 6.3. Thermoelectric properties of electronically conducting polymers Flexible thermoelectric devices require a special class of polymeric materials known as ‘‘electronically conducting polymers.” These are conjugated polymers with good electronic conductivity. Inorganic materials are classic thermoelectric materials, but they are unsuitable for use in low-temperature flexible thermoelectric devices. Due to their unique properties, electronically conducting polymers are in great demand for the research and development of flexible devices in various sectors. Examples of electronically conducting polymers used include polypyrrole (PPY), PANI, polythiophene (PTH), poly(3,4-ethylene dioxythiophene) (PEDOT), polyacetylene (PA), and their derivatives. These polymers demonstrate semiconductor characteristics with a variety of electronic band structures. Among these, poly(2,7-carbazole), PEDOT, and PEDOT:polystyrenesulfonate (PEDOT:PSS) have emerged as significant candidates for thermoelectric applications. Their merits include easy processability, good electronic conductivity, low thermal conductivity, low cost, tenability in various sizes and shapes, and environmentfriendly nature. The thermoelectric properties of electronically conducting polymers are briefly discussed in this section. 6.3.1. Poly(3,4-ethylenedioxythiophene) PEDOT is a derivative of the PTH family and one of the more significant electronically conducting polymers due to its low thermal conductivity, low density, easy handling, and environmental stability. It is an example of an organic semiconductor, which is used in a variety of applications such as thin-film transistors, solar cells, and supercapacitors. The chemical structure of PEDOT is depicted in Fig. 27. It possesses a simple electronic band structure with a bandgap of 0.9 eV. The Seebeck coefficient of PEDOT is 100 lV/K and 140 lV/K for p-type and n-type, respectively. The major drawback of using this polymer in thermoelectric devices is its low electronic conductivity [222]. Table 30 Thermoelectric properties of Mg1Ag1Sb1 [221]. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Mg1Ag1Sb1 Mg1Ag1Sb1 Mg1Ag1Sb1

13 27 67

122 137 152

1000 833 714

14.88 15.63 16.49

1.32 1.23 1.13

0.3 0.38 0.49

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Fig. 27. Chemical structure of PEDOT.

An n-type PEDOT exhibits a low electronic conductivity of 0.64  103 S/cm but a high Seebeck coefficient of 4008 lV/ K at room temperature [223]. The electronic conductivity of electronically conducting polymers can be increased by incorporating conducting species. PEDOT:tosylate (PEDOT:Tos), which is prepared directly by mixing the EDOT monomers and an oxidative solution of iron (III) tris-p-toluenesulfonate, exhibits a ZT value of 0.25 [224]. The electronic conductivity of PEDOT:Tos is found to increase with the interaction between the adjacent polymer chains and the Tos counterions on the outside of the chain stacks. Each Tos counterion helps balance the charges in the PEDOT chains. The oxidation levels in PEDOT affect thermoelectric properties such as electronic conductivity and Seebeck coefficient. In the case of electronically conducting polymers, conduction occurs via phonon-assisted hopping mechanisms. When the mass of the PEDOT:Tos sample is reduced, the carrier concentration decreases and the electronic conductivity fluctuates. In the case of highly oxidized PEDOT:Tos, an increase in electronic conductivity along with a decrease in Seebeck coefficient is observed [224]. 6.3.2. PEDOT:PSS Additives such as PSS in PEDOT are reported to increase the electronic conductivity and in turn enhance the thermoelectric properties. Addition of PSS to PEDOT leads to an increase in electronic conductivity of 9% compared with PEDOT:Tos. This increase in electronic conductivity of this organic semiconductor upon doping yields a ZT of 0.42 when dimethylsulfoxide (DMSO)-mixed PEDOT:PSS is treated with EG [66]. The entrapment of PSS in the PEDOT chain creates a conduction path for positive charges, as shown in Fig. 28. Recent studies have shown that conducting additives must be incorporated to increase the electronic conductivity of PEDOT:PSS films [224,225]. Addition of DMSO, N,N-dimethyl formamide, and tetrahydrofuran is known to enhance carrier transport and help to align the PEDOT chains [222,224]. Kong et al. synthesized PEDOT:PSS films by doping with urea, resulted in increased electronic conductivity and Seebeck coefficient, from 8.16 to 63.13 S/cm at 27 °C and from 14.5 to 20.7 lV/K, respectively[226]. Kim et al. developed a spin-coated, EG-treated, and DMSO-mixed PEDOT:PSS with a ZT value of 0.42 [66]. The doping process produces an isotropic effect on the thermal transport, which reduces the thermal conductivity. Further, the reduced dopant volume is found to optimize the carrier concentration and thereby increase the ZT value [66]. The thermoelectric performances of various PEDOT:PSS samples are summarized in Table 31. 6.3.3. Poly(2,7-carbazole) Polycarbazoles are unique materials with a high Seebeck coefficient due to charge localization, in the presence of oxidized nitrogen. This polymer exhibits a low electronic conductivity possibly arising from the charge-carrier pinning. The chemical

Fig. 28. Chemical structure of PEDOT:PSS.

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C. Gayner, K.K. Kar / Progress in Materials Science 83 (2016) 330–382 Table 31 Thermoelectric performance of PEDOT:PSS samples. Material

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/cm K2)

Thermal conductivity (W/m K)

ZT

Ref.

PEDOT nanotubes PEDOT:PSS + 5% DMSO PEDOT:Tos films EG-treated DMSO-mixed PEDOT:PSS

27 27 27 20

4008 13.5 220 72

0.64  103 570 67 880

0.010 0.104 3.24 4.6

– 0.34 0.37 0.24

– 0.0092 0.25 0.42

[223] [225] [224] [66]

Fig. 29. Chemical structure of poly(2,7-carbazole).

structure of poly(2,7-carbazole) presented in Fig. 29 indicates the presence of oxidized nitrogen, which can cause a decrease in mobility. Aich et al. synthesized doped poly(2,7-carbazole) and found that the dopant introduces a secondary alkyl chain to the carbazole along the conjugated backbone [227]. Doping in poly(2,7-carbazole) results in increased electrical conductivities up to 500 S/cm and a Seebeck coefficient of 70 lV/K. This polymer exhibits better stability in air, with a power factor of 19 lW/m K2 [227]. 6.3.4. Polyaniline PANI is also a promising material for use in next-generation thermoelectric devices due to its low cost and easy processing. The electronic conductivity of PANI is found to be dependent on the processing conditions and the disorder in the structure [228–230]. The chemical structure of PANI is depicted in Fig. 30. The three different structures observed in PANI are leucoemeraldine (y = 1), emeraldine (y = 0.5), and pernigraniline (y = 0), as shown in Fig. 30. PANI has a very poor thermoelectric response, but its use in flexible and low-cost thermoelectric devices has increased its demand. The electronic conductivity of PANI depends on the oxidation level, molecular arrangement, interchain separation, degree of crystallinity, and doping. Therefore, the electronic conductivity and Seebeck coefficient of PANI can be further increased via doping [231,232]. In HCl-doped PANI, the electronic conductivity and ZT are found to increase with increasing HCl concentration, accompanied by a decrease in the Seebeck coefficient. A ZT value of 2.67  104 at 430 K is obtained for 1 M HCl-doped PANI [232]. Sun et al. synthesized b-naphthalene sulfonic acid-doped PANI nanotubes with a Seebeck coefficient of 150–225 lV/K [233]. The disorder induced by b-naphthalene sulfonic acid interferes with electronic conduction, resulting in a low electronic conductivity of 7.7  103 S/cm [233]. Strategies for preparing nanocomposites of inorganic material/electronically conducting polymer can be explored further as they are promising candidates for use in flexible high-performance thermoelectric devices. 6.4. Thermoelectric properties of carbon nanomaterials Carbon is found in the form of a variety of allotropes, each of them exhibiting unique properties. Many allotropes of carbon exhibit high electrical and thermal conductivities, but controlling these parameters particularly for thermoelectric applications is challenging. The major drawback of using carbon materials for thermoelectric purposes is their high thermal conductivity. The ability of a material to conduct heat is based on its atomic arrangement. However, the thermal conductivity can be altered by manufacturing nanostructured materials. Nanowires do not conduct heat due to phonon–boundary scattering or phonon dispersions. The heat is transported in the solid material by acoustic phonons produced by the vibrations of

Fig. 30. Chemical structure of PANI.

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Fig. 31. The thermal conductivities of various carbon materials. Reproduced with permission from [235].

the crystal lattices as well as by electrons. The high thermal conductivity of carbon materials is the major limitation of their thermoelectric applications. Their strong covalent sp2 bonding (except for diamond, which has sp3 bonding) results in heat transfer via phonons produced by lattice vibrations [234,235]. The thermal conductivity of carbon materials can be significantly modified using their nanostructures such as CNTs, carbon nanowires, and graphene. Among these carbon nanomaterials, CNTs have gained much significance for their unique properties such as large surface area, good electrical conductivity, and high thermal stability. [236]. For example, experimental studies have shown that the thermal conductivity of multiwalled carbon nanotubes (MWCNTs) increases with decreasing number of walls because of the interaction of phonons and electrons at the multiple walls. A decrease in thermal conductivity from 2800 to 500 W/m K is noted for MWCNTs, when the outer diameter is increased from 10 to 28 nm [237]. The range of thermal conductivities of various carbon materials is presented in Fig. 31. The feasibility of using carbon materials in thermoelectric devices is still a matter of debate. 6.4.1. Carbon nanotubes CNTs are one-dimensional carbon nanomaterials with nanoscale diameters and single or multiple walls. They are considered potential candidates for thermoelectric devices. Doping in MWCNTs is found to alter the charge-carrier concentration. For example, Kunadian et al. synthesized boron- and nitrogen-doped MWCNTs, which exhibited p-type and n-type behavior, respectively [238]. Doping results in a decrease in thermal conductivity up to 75% arising from the induced defects inside the MWCNTs. The Seebeck coefficient of MWCNTs can be increased by increasing the density of charge carriers via oxygen doping. An increase in the doping concentration leads to an increase in the Seebeck coefficient. A thermoelectric module comprising five-cell p- and n-type MWCNTs exhibits a thermoelectric power of 16 lW at 27 °C [238]. Qin et al. used SPS to alter the orientations of MWCNTs along the direction of applied pressure without damaging their morphologies. Thus, they successfully obtained reduced thermal and electrical conductivities [239]. The operating temperature of MWCNTs for thermoelectric application can be varied from 87 to 567 °C, as shown in Fig. 32 [239]. Kim et al. developed a p- and n-type CNT film module composed of 172 thermocouples, with a thermopower of 465 mV at 49 °C [240]. The scanning electron microscope (SEM) images of individual p- and n-type thermoelements used in the module are presented in Fig. 33. The doping process determines the nanotube diameter; for example, n-type CNTs have larger diameter than p-type CNTs. Moreover, the nanotube diameter is found to affect electrical conduction as well [240]. 6.4.2. Graphene In thermoelectric research, the use of graphene as a thermoelectric material is highly debatable because of two major drawbacks: (i) j is very high and (ii) the Seebeck coefficient is very small due to the zero bandgap. However, studies have reported an electron mobility ranging from 1000 to 7000 cm2 V/s and a thermoelectric power of 80 lV/K at room temperature [241]. The sign of the thermoelectric power, which defines the majority of charge carriers, changes from positive to negative as the gate bias crosses the charge neutrality point [241,242]. Theoretical studies on zigzag graphene nanoribbons have shown that a ZT of 4 can be obtained at room temperature, based on the assumption that the lattice thermal conductivity is significantly reduced via phonon-edge disorder scattering without any effect on electron transport [243]. Hence, graphene is a promising thermoelectric material due to its high thermopower compared with elemental semiconductors [241,243]. Graphene with a lattice disorder intentionally introduced through electron beam irradiation or with charge impurities may exhibit a good thermoelectric response [244,245]. Disordered graphene can be used to enhance the thermoelectric performance of graphene. A thermopower of 50–100 lV/K at room temperature has been reported for the graphene sample exfoliated onto a thin layer of SiO2 [242,246]. Zuev et al. studied the variations in the Seebeck coefficient and electrical

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Fig. 32. FESEM image of sintered MWCNTs and its temperature-dependent electrical and thermal conductivities. Here, the notations \ and || denote the pressure orientation along the perpendicular and parallel directions, respectively. Reproduced with permission from [239].

Fig. 33. Cold fracture cross section of p-type (a) and n-type CNT film (b) [Inset: corresponding high-resolution image, scale bar = 100 nm]; surface morphology of p-type (c) and n-type (d) CNT film. Reproduced with permission from [240].

conductivity as a function of gate voltage (Vg). They found that the Seebeck coefficient of graphene is dependent on temperature and gate voltage, as well as on carrier density, as shown in Fig. 34 [241]. Sevincli et al. used a bottom-up fabrication route to increase the ZT value of graphene nanoribbons to as much as 3.25 at 800 K. Hybrid nanostructuring can decreased the thermal conductivity by up to 98.8% at room temperature. Thus, graphene nanoribbons synthesized with atomic

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Fig. 34. Plot of (a) electrical conductivity [Inset: SEM image of the typical thermoelectric device] and (b) thermoelectric power of a graphene as a function of Vg for T = 300, 150, 80, 40, and 10 K. Reproduced with permission from [241].

Fig. 35. Plot of variation in ZT of graphene nanoribbons as a function of length. Reproduced with permission from [247].

precision have clearly enhanced thermoelectric performance [247]. Fig. 35 plots the variations in ZT with respect to the length of graphene nanoribbons at different temperatures. 6.5. Thermoelectric properties of carbon nanomaterial/polymer nanocomposites Nanocomposites composed of CNTs or graphene as filler materials and electronically conducting (or, non-conducting) polymers as matrices have been developed recently. They have also been found to be suitable candidates for thermoelectric devices. Examples of such nanocomposites include PANI/graphene nanocomposites and PEDOT:PSS/graphene nanocomposites. Carbon nanomaterial/polymer nanocomposites can be used to enhance the thermoelectric properties and their flexibility in application. In a carbon nanomaterial/polymer nanocomposite, the effect of filler materials (e.g., CNTs, graphene, etc.) in the polymer matrices can be analyzed based on the ‘‘thermal conductivity enhancement factor (d).” It is defined as follows:

d¼

keff  kbase kbase

ð12Þ

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Fig. 36. (a) Plot of variation in thermopower and (b) electrical conductivity for layered MWCNT/PVDF nanocomposite film. Reproduced with permission from [250].

Fig. 37. Surface morphology of p–n junction of MWCNT/PVDF nanocomposite. Reproduced with permission from [250].

Fig. 38. (a–d) Schematic of p/n type of thermoelements, SEM image of thermoelement, p–n junction for fabricating a device, length scale of thermoelements and (e) power output of MWCNT-based device (Inserted image is an MWCNT-based module). Reproduced with permission from [251].

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Fig. 39. (a and b) TEM images of PEDOT:PSS/graphene films (scale bar: 500 nm); plot of variation in electrical conductivity and Seebeck coefficient (c), and power factor (d) as a function of graphene concentration. Reproduced with permission from [252].

Fig. 40. (a and b) TEM images of PANI/graphene nanocomposite and (c and d) digital photographs of flexible nanocomposite film. Reproduced with permission from [253].

where keff is the thermal conductivity of the nanocomposite and kbase is the thermal conductivity of the initial base material (i.e., the matrix) [235]. These emerging flexible nanocomposites can be used in novel thermoelectric technologies at room temperature.

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6.5.1. CNT/electronically conducting polymer nanocomposite Yao et al. synthesized a hybrid nanocomposite composed of single-walled carbon nanotubes (SWCNTs) and PANI with a Seebeck coefficient of 40 lV/K and an electrical conductivity of 125 S/cm [248]. With increase in SWCNT content in the PANI matrix, both the Seebeck coefficient and the electrical conductivity increase [248]. The electrical conductivity of a nanocomposite thin film composed of SWCNTs and polyvinylidene difluoride (PVDF) increases with an increase in PVDF concentration, although the reverse is observed for the Seebeck coefficient [249]. Hewitt et al. prepared a nanocomposite film consisting of MWCNTs and PVDF, which are layered into multiple element modules (denoted as ‘‘felt fabric”); these multiple layers are the sum of individual MWCNT/PVDF layers [250]. The power output of this module is high, due to the number and contribution of layers. The output voltage of the felt fabric increases with an increase in the number of fabric layers, with a reported power output of 235 lV/K at 27 °C for 11 alternate layers. In the near future, these nanocomposites may well replace the classical Bi2Te3, particularly for room temperature thermoelectric applications. The variations in thermopower and electrical conductivity for different layers of MWCNT/PVDF nanocomposite film at different temperatures are plotted in Fig. 36a and b, respectively. Fig. 37 represents the surface morphology of the p–n junction of the MWCNT/PVDF nanocomposite, which indicates that PVDF is uniformly coated over the surface of MWCNTs. The p- and n-type MWCNT/PVDF nanocomposites are joined by heating the polymer junction, such that only the polymer is melted and MWCNTs are visible at the junction, which participate further in the conduction process [250,251]. A maximum Seebeck coefficient of 96 lV/K at 27 °C is obtained for polyethylenimine-doped SWCNTs [251]. The multi-module device composed of nine modules is connected in series yielding a power output of 0.7 lW, the highest value reported for CNT-based polymer nanocomposites to date. The fabricated multi-module device and its performance are shown in Fig. 38 [251].

Fig. 41. Schematic showing the charge transport in the rGO–polymer and C60/rGO nanohybrid-filled polymer nanocomposite. Reproduced with permission from [254].

Fig. 42. Thermal conductivities and ZTs of nanohybrid-filled polymer nanocomposites. Reproduced with permission from [254].

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Table 32 Thermoelectric performance of various graphene/polymer nanocomposites. Material

Synthesis route

Operating temperature (°C)

Seebeck coefficient (lV/K)

Electrical conductivity (S/cm)

Power factor (lW/m K2)

Ref.

Pristine PEDOT:PSS PEDOT:PSS/graphene PEDOT:PSS/rGO PANI/graphene PANI/graphene PEDOT:PSS/rGO

Polymerization + mixing Mixing + spin coating Polymerization + mixing Mechanical grinding + cold pressing Polymerization + solution process Solution process + mixing

27 27 27 27 27 27

23.10 59 26.77 34 26 21

453 32.13 637 123 814 715

24.17 11.09 45.67 14 55 31.2

[252] [255] [252] [256] [253] [254]

6.5.2. Graphene/electronically conducting polymer composite Despite certain unresolved issues (such as the use of graphene as a thermoelectric material), research into the development of novel nanocomposites consisting of graphene and polymers is under way. The thermoelectric properties of graphene/polymer nanocomposites have been systematically studied as a function of filler content and at different operating temperatures. Yoo et al. prepared a PEDOT:PSS/graphene nanocomposite via an in situ polymerization method, and the surface topography of the nanocomposite is depicted in Fig. 39a and b. In situ polymerization produces a uniform coating of the polymer on the surface of graphene nanosheets, and the nanocomposite contains conducting channels throughout. This results in a high electrical conductivity of 637 S/cm and a Seebeck coefficient of 26.77 lV/K (Fig. 39c) [252]. A power factor of 45.67 lW/m K2 is obtained for the PEDOT:PSS/graphene nanocomposite, and the variation in power factor with respect to the graphene concentration in the nanocomposite is plotted in Fig. 39d. Wang et al. synthesized a PANI/graphene nanocomposite film, with a maximum power factor of 55 lW/m K2 for the nanocomposite [253]. The thermoelectric performance of the PANI/graphene nanocomposite film is enhanced by improving the molecular ordering of the composite film. This nanocomposite film exhibits an electrical conductivity of 814 S/cm with a Seebeck coefficient of 26 lV/K. The dispersion of graphene nanosheets without restacking introduces more nanointerfaces, eventually strengthening the p–p conjugation interactions between PANI and graphene nanosheets [253]. Fig. 40 represents the TEM images of the PANI/graphene nanocomposite, which indicates the uniform coating of PANI on the graphene nanosheets. Zhang et al. synthesized hierarchical nanohybrid-filled polymer nanocomposites composed of C60/rGO–PEDOT:PSS with a ZT of 0.067, which is one order of magnitude higher than the rGO–PEDOT:PSS nanocomposite [254]. In the nanohybridfilled polymer nanocomposite, fillers such as reduced graphene oxide (rGO) and fullerene (C60), and the polymer PEDOT: PSS are used. The nanohybrid-filled polymer nanocomposite showed enhanced charge transport due to the potential interfacial energy-filtering process. The charge transport mechanism in the nanohybrid-filled polymer nanocomposite is schematically presented in Fig. 41. The variation in thermal conductivity and ZT with respect to the rGO concentration in this nanocomposite is plotted in Fig. 42. This figure indicates that an initial increase in rGO concentration leads to an increase in both thermal conductivity and ZT, although with 30 wt% rGO in the nanocomposite, the thermal conductivity increases with a consequent decrease in ZT [254]. The thermoelectric response of various graphene/polymer nanocomposites are summarized in Table 32. 7. Perspective and concluding remarks In this review, the syntheses and thermoelectric properties of various inorganic materials are discussed in detail using examples. The syntheses and thermoelectric performances of certain emerging thermoelectric materials such as electronically conducting polymers, carbon nanomaterials, and their nanocomposites are also illustrated briefly. The various factors determining the thermoelectric properties of materials are elucidated with theoretical and experimental support. Promising approaches for enhancing the thermoelectric response of materials, such as alloying, use of nanocomposites and nanoinclusions, and doping, are described with examples. Electronically conducting polymers are highly flexible in their thermoelectric applications, although they exhibit poor thermoelectric responses. The thermoelectric response of electronically conducting polymers can be enhanced with nanocomposite strategies by preparing carbon nanomaterial/electronically conducting polymer nanocomposites. Acknowledgment The authors acknowledge the financial support provided by the Department of Science and Technology (DST), India, for conducted this research work. The authors thank Prof. Malay K. Das and Mr. Jayesh Cherusseri for their valuable inputs. References [1] Rowe DM. CRC handbook of thermoelectrics, macro to nano. CRC Press, Taylor & Francis Group; 2006. [2] Rosi FD. Thermoelectricity and thermoelectric power generation. Solid State Electron 1968;11:833–68. [3] Chasmar RP, Stratton R. The thermoelectric figure of merit and its relation to thermoelectric generators. J Electron Control 1959;7:52–72.

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